light metal systems. part 4: selected systems from al-si-ti to ni-si-ti

XI

Landolt-BörnsteinNew Series IV/11A4

MSIT®

Introduction

Introduction

Data Covered

The series focuses on light metal ternary systems and includes phase equilibria of importance for alloydevelopment, processing or application, reporting on selected ternary systems of importance to industriallight alloy development and systems which gained otherwise scientific interest in the recent years.

General

The series provides consistent phase diagram descriptions for individual ternary systems. Therepresentation of the equilibria of ternary systems as a function of temperature results in spacial diagramswhose sections and projections are generally published in the literature. Phase equilibria are described interms of liquidus, solidus and solvus projections, isothermal and pseudobinary sections; data on invariantequilibria are generally given in the form of tables.

The world literature is thoroughly and systematically searched back to the year 1900. Then, thepublished data are critically evaluated by experts in materials science and reviewed. Conflicting informationis commented upon and errors and inconsistencies removed wherever possible. It considers those, and onlythose data, which are firmly established, comments on questionable findings and justifies re-interpretationsmade by the authors of the evaluation reports.

In general, the approach used to discuss the phase relationships is to consider changes in state and phasereactions which occur with decreasing temperature. This has influenced the terminology employed and isreflected in the tables and the reaction schemes presented.

The system reports present concise descriptions and hence do not repeat in the text facts which canclearly be read from the diagrams. For most purposes the use of the compendium is expected to be self-sufficient. However, a detailed bibliography of all cited references is given to enable original sources ofinformation to be studied if required.

Structure of a System Report

The constitutional description of an alloy system consists of text and a table/diagram section which areseparated by the bibliography referring to the original literature (see Fig. 1). The tables and diagrams carrythe essential constitutional information and are commented on in the text if necessary.

Where published data allow, the following sections are provided in each report:

Literature Data

The opening text reviews briefly the status of knowledge published on the system and outlines theexperimental methods that have been applied. Furthermore, attention may be drawn to questions which arestill open or to cases where conclusions from the evaluation work modified the published phase diagram.

Binary Systems

Where binary systems are accepted from standard compilations reference is made to these compilations.In other cases the accepted binary phase diagrams are reproduced for the convenience of the reader. Theselection of the binary systems used as a basis for the evaluation of the ternary system was at the discretionof the assessor.

XII


MSIT®

Introduction

Solid Phases

The tabular listing of solid phases incorporates knowledge of the phases which is necessary or helpfulfor understanding the text and diagrams. Throughout a system report a unique phase name and abbreviationis allocated to each phase.

Phases with the same formulae but different space lattices (e.g. allotropic transformation) aredistinguished by:

– small letters (h), high temperature modification (h2 > h1)(r), room temperature modification(1), low temperature modification (l1 > l2)

– Greek letters, e.g., , '– Roman numerals, e.g., (I) and (II) for different pressure modifications.In the table “Solid Phases” ternary phases are denoted by * and different phases are separated by

horizontal lines.

Heading

Literature Data

Binary Systems

Solid Phases

Pseudobinary Systems

Invariant Equilibria

Liquidus, Solidus, Solvus Surfaces

Isothermal Sections

Miscellaneous

Miscellaneous

Isothermal Sections




Solid Phases

Binary Systems

Text

References

Tables anddiagrams

Temperature-Composition Sections

Temperature-Composition Sections

Thermodynamics

Notes on Materials Properties and Applications

Thermodynamics


Fig. 1: Structure of a system report

XIII


MSIT®

Introduction


Pseudobinary (quasibinary) sections describe equilibria and can be read in the same way as binary diagrams.The notation used in pseudobinary systems is the same as that of vertical sections, which are reported under“Temperature – Composition Sections”.


The invariant equilibria of a system are listed in the table “Invariant Equilibria” and, where possible, aredescribed by a constitutional “Reaction Scheme” (Fig. 2).

The sequential numbering of invariant equilibria increases with decreasing temperature, one numberingfor all binaries together and one for the ternary system.

Equilibria notations are used to indicate the reactions by which phases will be– decomposed (e- and E-type reactions)– formed (p- and P-type reactions)– transformed (U-type reactions)For transition reactions the letter U (Übergangsreaktion) is used in order to reserve the letter T to denote

temperature. The letters d and D indicate degenerate equilibria which do not allow a distinction accordingto the above classes.


The phase equilibria are commonly shown in triangular coordinates which allow a reading of theconcentration of the constituents in at.%. In some cases mass% scaling is used for better data readability(see Figs. 3 and 4).

In the polythermal projection of the liquidus surface, monovariant liquidus grooves separate phaseregions of primary crystallization and, where available, isothermal lines contour the liquidus surface (seeFig. 3).

Isothermal Sections

Phase equilibria at constant temperatures are plotted in the form of isothermal sections (see Fig. 4).

Temperature – Composition Sections

Non-pseudobinary T-x sections (or vertical sections, isopleths, polythermal sections) show the phasefields where generally the tie lines are not in the same plane as the section. The notation employed for thelatter (see Fig. 5) is the same as that used for binary and pseudobinary phase diagrams.

Thermodynamics

Experimental ternary data are reported in some system reports and reference to thermodynamicmodelling is made.


Noteworthy physical and chemical materials properties and application areas are briefly reported if theywere given in the original constitutional and phase diagram literature.

Miscellaneous

In this section noteworthy features are reported which are not described in preceding paragraphs. Theseinclude graphical data not covered by the general report format, such as lattice spacing – composition data,p-T-x diagrams, etc.

XIV


MSIT®

Introduction

Fig

ure

2:

T

ypic

al r

eact

ion s

chem

e

Ag-T

lT

l-B

iB

i-A

gA

g-T

l-B

i

(Tl)

(h)

(T

l)(r

),(A

g)

23

4d

1l

(A

g)

+ (

Bi)

26

1e 5

(Ag

) +

(T

l)(h

) +

Tl 3

Bi

L +

Tl 3

Bi

(A

g)

+ (

Tl)

(h)

28

9U

1

l (

Ag

)+(T

l)(h

)

29

1e 3

l (

Tl)

(h)+

Tl 3

Bi

30

3e 1

l (

Bi)

+T

l 2B

i 3

20

2e 7

l T

l 3B

i+T

l 2B

i 3

19

2e 8

(Tl)

(h)

Tl 3

Bi+

(Tl)

(r)

14

4e 9

L (

Ag

) +

Tl 3

Bi

29

4e 2

(max

)

L (

Ag

) +

(T

l)(h

)

28

9e 4

(min

)

L (

Ag

) +

Tl 2

Bi 3

20

7e 6

(max

)

(Ag

)+(B

i)+

Tl 2

Bi 3

L (

Ag)+

(Bi)

+T

l 2B

i 31

97

E1

(Ag

)+(T

l)(r

)+T

l 3B

i

(Tl)

(h)

Tl 3

Bi

+ (

Tl)

(r),

(Ag

)1

44

D1

(Ag

)+T

l 3B

i+T

l 2B

i 3

L (

Ag

)+T

l 3B

i+T

l 2B

i 31

88

E2

seco

nd b

inar

y

eute

ctic

rea

ctio

nfi

rst

bin

ary e

ute

ctic

rea

ctio

n

(hig

hes

t te

mper

ature

)te

rnar

y m

axim

um

reac

tion

tem

per

atu

re

of

26

1°C

mo

no

var

iant

equil

ibri

um

sta

ble

do

wn

to

lo

w

tem

per

ature

s

seco

nd

tern

ary

eute

ctic

reac

tion

equat

ion o

f eu

tect

oid

reac

tion a

t 144°C

XV


MSIT®

Introduction

20

40

60

80

20 40 60 80

20

40

60

80

A B

C Data / Grid: at.%

Axes: at.%

δ700

p1

500

400

400°C

γ

300

U e1

700

500

β(h)

400

300

E

300

α400

e2

500°C isotherm, temperature is usualy in °C

liquidus groove to decreasing temperatures

estimated 400°C isotherm

limit of known region

ternary invariantreaction

binary invariantreaction

primary γ-crystallization

20

40

60

80

20 40 60 80

20

40

60

80

Al B

C Data / Grid: mass%

Axes: mass%

L+γ

γ+β(h)

L+γ+β(h)

β(h)

L+β(h)

L

L+α

α

phase field notation

estimated phase boundary

tie line

three phase field (partially estimated)

experimental points(occasionally reported)

limit of known region

phase boundary

γ

Fig. 3: Hypothetical liquidus surface showing notation employed

Fig. 4: Hypothetical isothermal section showing notation employed

XVI


MSIT®

Introduction

References

The publications which form the bases of the assessments are listed in the following manner:[1974Hay] Hayashi, M., Azakami, T., Kamed, M., “Effects of Third Elements on the Activity of Lead

in Liquid Copper Base Alloys” (in Japanese), Nippon Kogyo Kaishi, 90, 51-56 (1974) (Experimental,Thermodyn., 16)

This paper, for example, whose title is given in English, is actually written in Japanese. It was publishedin 1974 on pages 51- 56, volume 90 of Nippon Kogyo Kaishi, the Journal of the Mining and MetallurgicalInstitute of Japan. It reports on experimental work that leads to thermodynamic data and it refers to 16 cross-references.

Additional conventions used in citing are:# to indicate the source of accepted phase diagrams* to indicate key papers that significantly contributed to the understanding of the system.Standard reference works given in the list “General References” are cited using their abbreviations and

are not included in the reference list of each individual system.

60 40 200

250

500

750

A 80.00B 0.00C 20.00

A 0.00B 80.00C 20.00Al, at.%

Tem

pera

ture

, °C

L

32.5%L+β(h)

β(r) - room temperature

β(r)

L+α+β(h)

α+β(h)

α

L+α

phase field notation

concentration ofabscissa element

alloy compositionin at.%

β(h)

modification

β(h) - high temperaturemodification188

temperature, °C

Fig. 5: Hypothetical vertical section showing notation employed

XVII


MSIT®

Introduction

General References

[C.A.] Chemical Abstarts - pathways to published research in the world's journal and patentliterature - http://www.cas.org/

[Curr.Cont.] Current Contents - bibliographic multidisciplinary current awareness Web resource - http://www.isinet.com/products/cap/ccc/

[E] Elliott, R.P., Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York(1965)

[G] Gmelin Handbook of Inorganic Chemistry, 8th ed., Springer-Verlag, Berlin [H] Hansen, M. and Anderko, K., Constitution of Binary Alloys, McGraw-Hill, New York

(1958) [L-B] Landolt-Boernstein, Numerical Data and Functional Relationships in Science and

Technology (New Series). Group 3 (Crystal and Solid State Physics), Vol. 6, Eckerlin, P.,Kandler, H. and Stegherr, A., Structure Data of Elements and Intermetallic Phases (1971);Vol. 7, Pies, W. and Weiss, A., Crystal Structure of Inorganic Compounds, Part c, KeyElements: N, P, As, Sb, Bi, C (1979); Group 4: Macroscopic and Technical Properties of

Matter, Vol. 5, Predel, B., Phase Equilibria, Crystallographic and Thermodynamic Data of

Binary Alloys, Subvol. a: Ac-Au ... Au-Zr (1991); Springer-Verlag, Berlin. [Mas] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, ASM, Metals Park, Ohio (1986) [Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International,

Metals Park, Ohio (1990) [P] Pearson, W.B., A Handbook of Lattice Spacings and Structures of Metals and Alloys,

Pergamon Press, New York, Vol. 1 (1958), Vol. 2 (1967) [S] Shunk, F.A., Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York

(1969) [V-C] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for

Intermetallic Phases, ASM, Metals Park, Ohio (1985) [V-C2] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for

Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

1


MSIT®

Al–Si–Ti

Aluminium – Silicon – Titanium

Pierre Perrot

Literature Data

The Ti-rich corner up to 13 at.% Al and 3 at.% Si was investigated from 600 to 1200°C [1954Tur, 1958Cro]

and partial isothermal sections were presented at 600, 800, 900, 1000°C [1954Tur, 1958Cro] and 1200°C

[1954Tur, 1962Sch]. 30 alloys were prepared by arc melting, annealed during various times (24 h at

1200°C, 720 h at 600°C) and examined by micrography. This investigation was further extended by

[1997Bul] which constructed the solidus and liquidus lines at 1300°C in the Ti-rich corner of the diagram,

gave the isopleth at 10 at.% Si. [1998Bul] showed a correspondence between the microhardness and the

solidus temperature in the same part of the diagram. [1999Aze] investigated phase relations in the Ti-rich

alloys with 14-22 at.% Al and 1-3.5 at.% Si by EDX analysis with construction the respective partial

isothermal section of the phase diagram at 700, 800, 900 and 1200°C. [2000Aze, 2002Aze] used the

experimental results [1999Aze] to calculate the partial isothermal sections of the Al-Si-Ti phase diagram at

700, 800, 900, 1000 and 1200°C. Calculated sections turned out to be in good agreement with experimental

data, although some discrepancies took place, yet. Ternary phases were identified mainly by [1961Bru,

1963Sch, 1965Ram]. The ternary phases and the solubility range from TiAl3 to Ti(Al0.8Si0.2)3 at 700°C,

were reported by [1965Ram]. They were obtained by X-ray analysis of 39 alloys after annealing for 2 h to

5 d at 700°C with the statement that equilibrium was not always obtained. [1968Kam] constructed an

isothermal section of the phase diagram at room temperature by microstructural and X-ray analysis and

found a ternary compound, in equilibrium with aluminium, assumed to be TiSi2Al. Its crystal structure and

exact chemical composition, however, were not established. [1978You] calculated the Al-rich corner of the

diagram and presented partial vertical sections at 0.05, 0.2 and 0.5 at.% Si. Owing to the lack of

thermodynamic data, the ternary compounds were not taken into account, leading to some discrepancies

with the Al-rich corner of the phase diagram constructed by [1984Ory] using lattice parameter

measurements. [1988Zak] studied phase equilibria from 550 to 850°C, 10 to 14 mass% Si and 0 to 0.6

mass% Ti by chemical analysis, DTA, metallographic and X-ray analysis. A ternary phase Ti2Al3Si2 was

observed in the system. This phase takes part in the invariant transition reaction L+Ti2Al3Si2 (Al)+(Si) at

579°C. Three new other invariant transition reactions with Ti2Al3Si2 participation in the Al and Si corners

of the system above 579°C were supposed, as well. According to [1992Zak], crystal structure of the

Ti2Al3Si2 compound differs from those of TiAl3 and TiSi2, but it was not identified. [1976Mon] suggested

that the phases Ti7Al5Si12, Ti2Al3Si2 as well as TiAlSi2 are solid solutions of Al in the TiSi2 phase.

However, [2002Sah] showed by X-ray diffraction analysis, that in rapid cooling condition the addition of

Ti to Al-17.5 mass% Si alloy leads to the formation of TiAl3 and TiAlSi2 whose structure is different from

that of both TiAl3 (D022 tetragonal) and TiSi2 (C54 orthorhombic). [1990Wu] used microscopy method for

determination of the eutectic Ti3(Al,Si)+Ti5(Si,Al) line position in limits of 5.0-18.75 at.% Si and 18.75-30

at.% Al. [1994Man] constructed approximate partial liquidus projection of the phase diagram in vicinity of

the TiAl phase using microstructure, chemical analysis, EDX analysis and calculations. [1994Wu]

determined by the special investigation position of the ( Ti)+Ti5Si3 eutectic line in the ternary phase

diagram up to 38.5 at.% Al. [1996Li, 1999Li] studied effect of Si addition on the ( 2)/( ( 2)+ ) phase

boundary in the Al-Ti binary system using electron probe microanalysis (EPMA) and thermodynamic

analysis. A shift of the boundary to the Ti-rich side was established. [1995Per] calculated a schematic

isothermal section of the Al-Si-Ti ternary phase diagram at 1100°C using only information on the binary

systems. The section constructed did not take into account solid solutions, neither ternary phases. This

diagram is qualitatively similar to the experimental ones.

Binary Systems

The binary systems Al-Si, Al-Ti and Si-Ti are respectively taken from [2003Luk], [2003Sch] and

[1987Mur] taking into account the thermodynamic optimization of [1996Sei]. The TiSi2 intermetallic

2


MSIT®

Al–Si–Ti

compound is dimorphous with the resultant structure depending on the method of synthesis. One of the

forms of TiSi2 [1956Cot] is metastable; however, it is stabilized by Al [1961Bru] leading to a continuous

solid solution from Ti(Si0.85Al0.15)2 to Ti(Si0.7Al0.3)2.

Following [1983Kub] and [1988Mur], the Ti9Al23 phase described by [1965Ram] is considered as

metastable. Actually, Ti9Al23 can be easily considered as an ordered, metastable modification of the Ti2Al5phase.

[1983Loi] reported the existence of three phases around the TiAl composition. Ti46Al54 and Ti42Al58

correspond to an ordered superstructure of the Cu type different from AuCu. The same authors [1988Loi]

consider TiAl2 as a reentrant phase with a structure type ZrGa2 above 1250°C and below 700°C, and a

structure type HfGa2 in between. The ZrGa2 form is likely metastable [2001Bra].

Solid Phases

Table 1 gives crystallographic data of all binary and ternary phases. The ternary phase Ti(AlxSi1-x)2 ( 2) has

been determined by two investigators [1961Bru, 1965Ram]. [1963Sch] detected by X-ray diffraction the

ternary phase Ti7Al5Si12 ( 1), stable below 900°C. The composition of another reported ternary phase, 3,

[1965Ram] was not identified. Two other ternary phases were revealed. One is Ti2AlSi3 (pseudotetragonal,

oC12, ZrSi2 type), identical to that of Ti(AlxSi1-x)2 given by [1961Bru] within the accuracy of the study.

[1965Ram] found a different ternary compound with the oC12, ZrSi2 crystal structure; however, its

composition could not be identified. Owing to the lack of experimental evidence for the ternary phases

TiAlSi2 [1968Kam] and Ti2Al3Si2 [1988Zak, 1992Zak], it is probable that these phases are identical to

Ti7Al5Si12 ( 1) and Ti(AlxSi1-x)2 ( 2), respectively.

Al appears to have little effect upon the solubility of Si in the ( Ti) phase and to decrease the solubility of

Si in the ( Ti) phase. The solubility of Si at 840°C decreases from 0.6 at.% in ( Ti) to 0.4 at.% in (Ti-Al12)

[1963Luz]. According to [1997Bul], the maximum solubility of Si in ,TiAl decreases from about 0.8 at.%

Si at the aluminium poor boundary to 0.6 at.% Si at the aluminium rich boundary. At 1200°C, the solubility

of Si in the phase decreases from 4 at.% Si for Al free ( Ti) to 1.75 at.% Si for (Ti-Al25) [1972Nar].

[1999Aze] agrees with the decreasing of the solubility of Si in ( Ti). However, they propose a high

solubility of Si in ( Ti) (about 1 at.%) which increases with the Ti content. This result contradicts the

observed hardening of Al-Ti alloys by precipitation of Ti5Al3 after addition of less than 1 at.% Si [1996You,

2000Bul, 2002Sun]. The ,TiAl phase dissolves less than 1 at.% Si [1976Sid]. At the solidus temperatures,

the ( Ti) homogeneity region stretches from 0 % Al and 4.7 at.% Si at 1330°C to 44.8 at.% Al and 0 at.%

Si at 1490°C. In the region in equilibrium with Ti5Si3, which extends from 48 to 51 at.% Al, the silicon

content is about 0.5 at.%. These experimental results [1997Bul] are not clearly confirmed by the Calphad

evaluation of [2000Aze, 2002Aze]. The maximum of aluminium solubility in Ti5Si3 is 9 at.% at ~1300°C

[1997Bul]. The homogeneity region extends towards the Al-Ti side between the 61 at.% and 65 at.% Ti

isopleths, which proves the substitution of silicon by aluminium [1997Bul]. The solubility of Ti in solid Al

decreases with the presence of Si [1985Guz].

Addition of Si to a Al-Ti alloy leads to a shift of the [ 2/( 2+ )] phase boundary to the Ti-rich side [1996Li,

1999Li]. The addition of 0.3 at.% Si increases the / transus by about 80-110°C with reference to the Al-Ti

binary system [1999Li].

[1941Pan] observed the precipitation of an additional compound when adding 0.6 at.% Ti to an Al-Si

eutectic alloy. [1957Now] confirmed its existence by annealing three alloys in the Al-TiSi2 section.

[1961Bru] identified a ternary compound as Ti(AlxSi1-x)2 (0.15 x 0.3) pseudotetragonal using X-ray

powder diffraction analysis. The same team by X-ray analysis, investigated 90 ternary alloys. They

observed a solubility of Si in TiAl3 of up to the composition Ti(Al0.85Si0.15)3.


An invariant transition reaction L+ 2 (Al)+(Si) at 579°C is reported by [1988Zak] with that temperature

being by 1.5 K above the binary Al-Si eutectic temperature measured as 577.5°C by the same authors. The

composition of “Ti2Al3Si2” was determined by chemical analysis of crystals separated by wet acid

dissolution of the (Al) phase. The X-ray diagram of “Ti2Al3Si2” is different from those of TiAl3 and TiSi2.

3


MSIT®

Al–Si–Ti

As the authors did not report the X-ray peaks and did not compare with the X-ray pattern of 1, one may

identify “Ti2Al3Si2” with 1 (Ti7Al5Si12), assuming the result of the chemical analysis was shifted by

undissolved inclusion of the (Al) phase. The invariant equilibria experimentally confirmed, mainly by

[1997Bul] are given in Table 2. The temperatures of the invariant points U1 and U2 (1420 and 1415°C)

accepted from [1997Bul] agree with the upper limit proposed by [1994Man], respectively T (U1) < 1480°C

and T (U2) < 1450°C. [1997Bul] presents the U2 invariant equilibrium as being of the E type

(L ( Ti)+ +Ti5Si3). Actually, the reaction is more likely of the U type (L+( Ti) +Ti5Si3) because the

coordinates of the invariant point fall outside of the ( Ti)+ +Ti5Si3 triangle. In addition to them two other

invariant reactions in solid state were shown by [1997Bul]. Both of them are included in Table 2 and the

reaction scheme (Fig. 3) as E1 an U3 equilibria. Unlike [1997Bul], the reaction at ~1035°C is assumed to

be of the eutectoid type instead of transition one because of connection with relevant three phase

monovariant reactions. The invariant U4 at 579°C was reported by [1988Zak] with a ternary compound

designated here as 1 taking part in the invariant transition reaction. It is more likely that 1 be the end of

the solid solution (Ti1-xAlx)8(AlySi1-y)16 with x = 0.12 and y = 0.25, that is approximately Ti7Al5Si12.

From the thermodynamic assessments of [1983Lia, 1988Mur, 1989Vah], an invariant equilibrium in the

solid state: Ti5Si3+Ti3Al Ti3Si+TiAl is calculated to occur at 1067°C.

Liquidus Surface

The liquidus surface near the Al-Si binary eutectic, determined by [1988Zak], is given in Fig. 1. A liquidus

surface of the whole Al corner was published by [1968Kam] but shows a univariant three-phase equilibrium

near 12 mass% Si going above 800°C and thus incompatible with [1988Zak]. It was therefore omitted here.

The eutectic valley near the Ti rich part of the diagram was first examined by [1990Wu], but the eutectic

line lies too close to the Al-Ti border and cannot be accepted. Following investigations were conducted by

[1993Zha, 1994Man, 1994Wu], then more precisely by [1997Bul, 1998Bul] which determined the tie lines

in the two phase domain ( Ti)+Ti5Si3 and confirmed the presence of a maximum e1 at 1545°C (1534°C

from [1994Wu]. The projections of the solidus and liquidus surfaces are presented in Fig. 2 and a partial

reaction scheme is given in Fig. 3.

Isothermal Sections

Partial isothermal section at 1523°C was calculated by [1994Man] from experimental isopleths considering

Ti5Si3 a stoichiometric binary compound. It is presented in Fig. 4 assuming existence of the homogeneity

range for Ti5Si3 taking into account the more realistic shape derived from the experimental work of

[1997Bul] at 1300°C. The Fig. 5 presents the phase equilibria at 1300°C [1997Bul]. In the Fig. 5 the

experimental tie lines in the two-phase domain ( Ti)+Ti5Si3 together with the experimental shape of the

Ti5Si3 single-phase domain are also reported. The isothermal section at 1200°C given in Fig. 6 is mainly

based on [1962Sch] with some changes concerning the existence of the ( Ti) phase from the Al-Ti binary,

which is also consistent with the isothermal section of [1954Tur, 1958Cro]. [2002Aze] presents somewhat

different isothermal sections of the Ti rich corner calculated from 700 to 1200°C which agree only

qualitatively with the Al-Ti binary system and do not take into account the high solubility (9 at.%) of Al

into Ti5Si3, so that experimental results seem to be more acceptable than calculated ones. [2002Aze] is

actually a reprint of [2000Aze] corrected from scaling errors made in the figures. Equilibrium between

phases TiAl and Ti5Si3 is confirmed by the experiments of [2001Boh, 2001Sun, 2002Sun] related to the

precipitation of Ti5Si3 in Si-bearing TiAl alloys and by those of [2002Hok] which prepared composite

structures of these two phases using explosive energy from underwater shock-waves. The Ti-Ti5Si3composite obtained by the same technique is explained by the low solubility of Si in TiAl3 at the low

temperature of the reaction. Figure 7 gives the isothermal section at 700°C, mainly from [1965Ram]. Minor

adjustments have been made to comply with the binary phase diagrams. The three-phase equilibrium

( Ti)+ 2+Ti3Si at 700°C is calculated from the thermodynamic assessments of [1983Lia, 1988Mur,

1989Vah]. Figure 8 giving the isothermal equilibria of the Ti corner at 1200 and 1000°C and partially at

800°C, is mainly based on [1954Tur] and [1958Cro]. However, the original diagrams have been modified

4


MSIT®

Al–Si–Ti

to include the Ti3Si phase which forms peritectoidically below 1170°C and to match with the Ti-Si and

Al-Ti accepted binary systems.


Several vertical sections were proposed: at Ti/Al = 3 [1972Nar], at 0.05, 0.2 and 0.5 at.% Si [1978You], at

10, 12 and 14 mass% Si up to 0.6 mass% Ti [1988Zak], and at 10 at.% Si [1997Bul]. Several isopleth for

xSi = 0.02, 0.035 and 0.05 were constructed using the ThermoCalc software [1994Man].

Thermodynamics

Thermodynamic assessments have been carried out for the three binary systems: Al-Si [1984Mur,

1986AnM, 2000Han], Al-Ti [1983Lia, 1988Mur, 1997Zha] and Si-Ti [1979Kau, 1989Vah, 1996Sei].

[1995Per] using only information on the binary systems, calculated a schematic Al-Si-Ti ternary diagram

at 1100°C, that is without taking into account solid solutions, neither ternary phases. This diagram is

qualitatively similar to the experimental one shown in Fig. 6 and explains the diffusion pathways observed

with diffusion couples 2-TiSi2, namely: 2- ,TiAl-TiAl2-Ti5Si4-TiAl3-TiSi-TiSi2, Integral enthalpies of

mixing are given for the Al-Si and Al-Ti binary liquids by [1987Des] at 1727°C and for the ternary liquid

at 1600°C near the Al-Si side by [1986Sud]. These data show strong attractive interactions between Ti and

Al as well as between Ti and Si. Calphad assessment of the Ti-rich part of the diagram (< 25 at.% Al and

< 5 at.% Si) has been carried out and isothermal sections has been drawn between 700 and 1200°C

[2000Aze, 2002Aze]. They are not included in this report because they contradict too much with

experimental data.


The high melting temperature of Ti5Si3 (2130°C) can be used for heat resistant materials and other

applications requiring high hardness, high oxidation resistance and high thermal conductivity. The synthesis

of TiAl-Ti5Si3 composites using explosion energy [2002Hok] produces materials with a hardness higher

than that of commercially available Ti5Si3 (HV 800-1050 kg mm-2 [1998Bul]). Using mechanical

alloying it is possible to prepare compounds Ti5(Si,Al)3 oversaturated in Al, in which 60 % of the Si atoms

are replaced by Al ones [1997Gua]. On another hand, structural applications of ,TiAl are of importance

due to its low density, high specific strength and stiffness at elevated temperatures [2001Sun]. Main

limitations of ,TiAl are poor ductility and toughness at room temperature. Since the solubility of Si in TiAl

is low, Si appears to be one of the most attractive candidates for raising the creep resistance of the alloy

[1997Gua]. Additions of 0.5 at.% Si to TiAl alloys result in formation of fine particles of Ti5Si3 precipitated

at / or / 2 grain boundaries, thus improving significantly mechanical properties [2000Rao, 2001Boh,

2001Sun, 2002Sun] of Al-Ti alloys.

Miscellaneous

[1997Via] used a simplified section of the Al-Si-Ti diagram at 1000°C drawn from [1962Sch, 1965Ram]

to obtain a representation of the phase equilibria in the Al-C-Si-Ti quaternary system with the aim of a better

understanding of the reaction processes likely to develop at the interface between silicon carbide

reinforcements and titanium aluminides matrices.

References

[1941Pan] Panserl, C., Guastalla, B., “Modification of Eutectic Al-Si Alloys. I. Influence of Ti

Additions as the Third Component” (in Italian), Alluminio, 10, 202-227 (1941) (Equi.

Diagram, Experimental, 161)

[1954Tur] Turney, D.H., Crossey, F.A., “Studies of Phase Relationships and Transformation Processes

of Ti Alloy System. Part VI: The Ti-Rich Corner of the Ti-Al-Si System”, Wright Air

5


MSIT®

Al–Si–Ti

Development Center, Technical Report, 54-101, 52-66 (1954) (Equi. Diagram,

Experimental, #, 26)

[1956Cot] Cottner, P.G., Kohn, J.A., Potter, R.A., “Physical and X-Ray Study of the Disilicides of Ti,

Zr and Hf”, J. Am. Ceram. Soc., 39, 11-12 (1956) (Crys. Structure, 8)

[1957Now] Nowotny, H., Huschka, H., “Studies of the Partial Systems Al-TiSi2, Al-ZrSi2, Al-MoSi2and Al-WSi2” (in German), Monatsh. Chem., 88, 494-501 (1957) (Crys. Structure, 21)

[1958Cro] Crossey, F.A., Turney, D.H., “Ti-Rich Corner of the Ti-Al-Si System”, Trans. Metall. Soc.

AIME, 212, 60-63 (1958) (Equi. Diagram, Experimental, #, 8)

[1961Bru] Brukl, C., Nowotny, H., Schob, O., Benesovsky, F., “The Crystal Structure of TiSi,

Ti(Al,Si)2 and Mo(Al,Si)2” (in German), Monatsh. Chem., 92, 781-788 (1961) (Crys.

Structure, 13)

[1962Sch] Schob, O., Nowotny, H., Benesovsky, F., “The Ternary Systems Formed by Ti, Zr and Hf

with Al and Si” (in German), Planseeber. Pulvermetall., 10, 65-71 (1962) (Equi. Diagram,

Experimental, #, *, 26)

[1963Luz] Luzhnikov, L.P., Novikhova, V.M., Marsev, L.P., “Solubility of b-Stabilizers in aTi” (in

Russian), Metall. Term. Obrab. Metallov, (2), 13-16 (1963) (English translation pp. 78-81)

(Experimental, 4)

[1963Sch] Schubert, K., Frank, K., Gohle, R., Maldonado, A., Meissner, H.G., Raman, A.,

Rossteutscher, W., “Structural Data on Metallic Phases VII” (in German),

Naturwissenschaften, 50, 41 (1963) (Crys. Structure, Review, 0)

[1965Ram] Raman, A., Schubert, K., “On the Constitution of Some Alloy Series Related to TiAl3. II.

Investigation on Some Systems T-Al-Si and T4...6-In” (in German), Z. Metallkd., 56, 44-52

(1965) (Crys. Structure, Equi. Diagram, #, *, 13)

[1968Kam] Kamei, K., Ninomiya, T., Hayashi, S., “The Phase Diagram of the Al-Si-Ti Equilibrium

System”, Met. Abstr. Light Metals and Alloys, 1964-1966, 67-70 (1968) (Equi. Diagram,

Experimental, #, 0)

[1972Nar] Nartova, T.T., Andreev, O.N., “Structure and Properties of Ti3Al and Alloys Based on it”

(in Russian), Stroenie, Svoitsva i Primenenie Metallov, 2nd Mater. Symp., 1972, Nauka,

Moscow, 194-197 (1974) (Equi. Diagram, Experimental, #, 14)

[1976Mon] Mondolfo, L.F., “Al-Si-Ti” Aluminium Alloys: Structure and Properties, Butterworth,

London, 614-615 (1976) (Review, 15)

[1976Sid] Sidorenko, F.A., Radovskii, I.Z., Chemerinskaya, L.S., Geld, P.V., “Structure and Magnetic

Susceptibility of Mutual V-Al and Ti-Si Solid Solutions with TiAl” (in Russian), Fiz.

Svoistva Met. Splavov, 1, 10-15 (1976) (Crys. Structure, Experimental, 11)

[1978You] Youdelis, W.V., “Calculated Al-Si-Ti Phase Diagram and Interpretation of Grain

Refinement Results”, Met. Sci., 12, 363-366 (1978) (Thermodyn., 12)

[1979Kau] Kaufman, L., “Coupled Phase Diagrams and Thermochemical Data for Transition Metals

Binary Systems -VI-“, Calphad, 3, 45-76 (1979) (Thermodyn., #, 16)

[1983Kub] Kubachewski, O., “Titanium. Physicochemical Properties of Its Compounds and Alloys”,

Atomic Energy Rev., Spec. Iss., No. 9, I.A.E.A. Vienna, 50-51 and 77-82 (1983) (Review,

Thermodyn.)

[1983Lia] Liang, W.W., “A Thermodynamical Assessment of the Al-Ti System”, Calphad, 7, 13-20

(1983) (Thermodyn., 27)

[1983Loi] Loiseau, A., Lasalmonie, A., “New Ordered Superstructure in non Stoichiometric TiAl” (in

French), Acta Crystallogr., B, 39, 580-587 (1983) (Crys. Structure, Experimental, 14)

[1984Mur] Murray, J.L., McAlister, A.J., “The Al-Si (Aluminum - Silicon) System”, Bull. Alloy Phase

Diagrams, 5, 74-84 (1984) (Thermodyn., #, *, 73)

[1984Ory] Orinbekov, S.B., Makanov, U.M., Guzei, L.S., Sokolovskaya, E.M., “The Interaction of Si

and Ti with Al” (in Russian), Vestn. Mosk. Univ., Ser. Khim., 25, 500-503 (1984) (Equi.


6


MSIT®

Al–Si–Ti

[1985Guz] Guzei, L.S., Kuznetsov, S.B., Orinbekov, S.B., Sokolovskaya, E.M., Makanov, U.M.,

"Phase Equilibria in the Al Corner of the System Al-Si-Cu-Ti" (in Russian), Vestn. Mosk.

Univ., Ser.2: Khim., 26, 393-395 (1985) (Experimental, 5)

[1986AnM] An Mey, S., Hack, K., “A Thermochemical Evaluation of the Si-Zn, Al-Si and Al-Si-Zn

Systems”, Z. Metallkd., 77, 454-459 (1986) (Thermodyn., #, 39)

[1986Sud] Sudavtsova, V.S., Batalin, G.I., Tutevitch, V.S., “Heat of Mixture of Liquid Alloys in the

Si-Al-Ti System” (in Russian), Ukr. Khim. Zh., 52, 1029-1031 (1986) (Experimental,

Thermodyn., 5)

[1987Des] Desai, P.D., “Thermodynamical Properties of Selected Binary Aluminum Alloys Systems,

Part 7: Al-Si, Part 8: Al-Ti”, J. Phys. Chem. Ref. Data, 16, 120-124 (1987) (Review,

Thermodyn., #, 29)

[1987Mur] Murray, J.L., “The Si-Ti (Silicon-Titanium) System”, Phase Diagrams of Binary Titanium

Alloys, ASM, Metals Park, OH 291-294 (1987) (Equi. Diagram, Crys. Sructure,

Thermodyn., Review, #, 29)

[1988Loi] Loiseau, A., Vannuffel, C., “TiAl2, a Reentrant Phase in Ti-Al System”, Phys. Status

Solidi, A, 107, 665-671 (1988) (Crys. Structure, 21)

[1988Mur] Murray, J.L., “Calculation of the Ti-Al Phase Diagram”, Metall. Trans. A, 19, 243-247

(1988) (Thermodyn., 23)

[1988Zak] Zakharov, A.M., Gildin, I.T., Arnold, A.A., Matsenko, Yu.A., “Phase Equilibria in the

Al-Ti-Si System at 10-14 % Si and 0-6 % Ti”, Russ. Metall., (4), 185-189 (1988), translated

from Izv. Akad. Nauk SSSR, Met., (4), 181-186 (1988) (Equi. Diagram, Experimental, #, 10)

[1989Vah] Vahlas, C., Chevalier, P.Y., Blanquet, E., “A Thermodynamic Evaluation of Four Si-M

(M = Mo, Ta, Ti, W) Binary Systems”, Calphad, 13(3), 273-292 (1989) (Thermodyn., #, 67)

[1990Sch] Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”,

Z. Metallkd., 81, 389-396 (1990) (Equi. Diagram, Experimental, 33)

[1990Wu] Wu, J.S., Beaven, P.A., Wagner, R., “The Ti3(Al, Si) + Ti5(Si,Al)3 Eutectic Reaction in the

Ti-Al-Si System”, Scr. Met. Mater., 24(1), 207-212 (1990) (Equi. Diagram,

Experimental, 5)

[1992Kat] Kattner, U.R., Lin, J.-C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the

Ti-Al System”, Metall. Trans. A, 23A(8), 2081-2090 (1992) (Equi. Diagam, Thermodyn.,

Assessment, *, #, 51)

[1992Zak] Zakharov, A.M., Guldin, I.T., Arnold, A.A., Matsenko, Yu.A., “Phase Equilibria in

Multicomponent Aluminum Systems with Copper, Iron, Silicon, Manganese and Titanium”

(in Russian), Metalloved. i Obrab. Tsv. Splavov: To 90th Ann. of Acad. A.A. Bochvar. RAN.

Int. Metallurgii, M, 6-17 (1992) (Equi. Diagram, 15)

[1993Zha] Zhang, L.T., Qiu, G.H., Wu, J.S., “Thermodynamic Calculation of the Ti5(Si, Al)3+ -Ti(Al,

Si) Eutectic Reaction”, Proceedings of the 7th National Symposium on Phase Diagrams,

206-210 (1993), Abstract in Red Book, MSI, 38(2), 868-869 (1993) (Equi. Diagram, #, 1)

[1994Man] Manesh, S.H., Flower, H.M., “Liquidus Projection of Ti-Al-Si Ternary System in Vicinity

of Alloys”, Mater. Sci. Technol., 10(8), 674-679 (1994) (Equi. Diagram, Experimental,

Thermodyn., #, 15)

[1994Wu] Wu, J.S., Qiu, G.H., Zhang, L.T., “The -Ti(Al,Si) + Ti5(Si,Al)3 Eutectic Line in the

Ti-Al-Si System”, Scr. Metall. Mater., 30(2), 213-218 (1994) (Equi. Diagram,

Experimental, #, 6)

[1995Per] Perepezko, J.H., da Silva Bassani, M.H., Park, J.S., Edelstein, A.S., Everett, R.K.,

“Diffusional Reactions in Composite Synthesis”, Mater. Sci. Eng. A, 195, 1-11 (1995)

(Calculation, Equi. Diagram, 41)

[1996Li] Li, J., Hao, S., “ ( )/ Phase Equilibria in Ti-Al-Si”, (in Chinese) Acta Metall. Sin.

(China), 32(11), 1171-1176 (1996) (Equi. Diagram, Experimental, 8)

[1996Sei] Seifert, H.J., Lukas, H.L., Petzow, G., “Thermodynamic Optimization of the Ti-Si System”,

Z. Metallkd., 87(1), 1-13 (1996) (Equi. Diagram, Thermodyn., Assessment, #, 63)

7


MSIT®

Al–Si–Ti

[1996You] You, B.-S., Park, W.-W., “Age Hardening Phenomena and Microstructure of Rapidly

Solidified Al-Ti-Si and Al-Cr-Y Alloys”, Scr. Mater., 34(2), 201-205 (1996) (Equi.

Diagram, Mechan. Prop., Experimental, 11)

[1997Bul] Bulanova, M., Tretyachenko, L., Golovkova, M., “Phase Equilibria in the Ti-rich Corner of

the Ti-Si-Al System”, Z. Metallkd., 88(3), 256-265 (1997) (Equi. Diagram, Experimental,

#, *, 15)

[1997Gua] Guan, Z.Q., Pfulmann, Th., Oehring, M., Bormann, R., “Phase Formation During Ball

Milling and Subsequent Thermal Decomposition of Ti-Al-Si Powder Blends”, J. Alloys

Compd., 252, 245-251 (1997) (Equi. Diagram, Experimental, 24)

[1997Via] Viala, J.C., Peillon, N., Bosselet, F., Bouix, J., “Phase Equilibria at 1000°C in the

Al-C-Si-Ti Quaternary System: an Experimental Approach”, Mater. Sci. Eng. A, A229,

95-113 (1997) (Equi. Diagram, Experimental, Thermodyn. 35)

[1997Zha] Zhang, F., Chen, S.L., Chang, Y.A., Kattner, U.R., “A Thermodynamic Description of the

Ti-Al System”, Intermetallics, 5, 471-482 (1997) (Equi. Diagram, Thermodyn.,

Assessment, #, 45)

[1998Bul] Bulanova, M., Soroka, A., Tretyachenko, L., Stakhov, D., “Microhardness of Structure

Units in the Ternary Ti-rich Ti-Si-Al Alloys”, Z. Metallkd., 89(6), 442-444 (1998) (Equi.


[1999Aze] Azevedo, C.R. de F., Flower, H.M., “Microstructure and Phase Relationships in Ti-Al-Si

System”, Mater. Sci. Technol., 15, 869-877 (1999) (Equi. Diagram, Experimental, 50)

[1999Li] Li, J., Zong, Y., Hao, Sh., “Effects of Alloy Elements (C, B, Fe, Si) on the Ti-Al Binary

Phase Diagram”, J. Mater. Sci. Technol., 15(1), 58-62 (1999) (Equi. Diagram,

Experimental, 13)

[2000Aze] Azevedo, C.R.F., Flower, H.M., “Calculated Ternary Diagram of Ti-Al-Si System”, Mater.

Sci. Technol., 16, 372-381 (2000) (Calculation, Equi. Diagram, Thermodyn., 15)

[2000Bul] Bulanova, M., Ban’kovsky, O., Soroka, A., Samelyuk, A., Tretyachenko, L., Kulak, L.,

Firstov, S., “Phase Composition, Structure and Properties of Cast Ti-Si-Sn-Al Alloys”,

Z. Metallkd., 91(1), 64-70 (2000) (Crys. Structure, Equi. Diagram, Experimental, Mechan.

Prop., 22)

[2000Han] Hansen, H.C., Lopper, C.R., “Effect of the Antimony on the Phase Equilibrium of Binary

Al-Si Allooys”, Calphad, 24, 339-352 (2000) (Equi. Diagram, Calculation, Thermodyn.,


[2000Rao] Rao, K.P., Du, Y.L., “In Situ Formation of Titanium Silicides Reinforced TiAl Based

Composites”, Mater. Sci. Eng. A, A277, 46-56 (2000) (Crys. Structure, Equi. Diagram,

Experimental, 38)

[2001Boh] Bohn, R., Fanta, G., Klassen, T., Bormann, R., “Mechanical Behaviour and Advanced

Processing of Nano- and Submicro-Grained Intermetallic Compound Based on -TiAl”,

Scr. Mater., 44(8-9), 1479-1482 (2001) (Equi. Diagram, Experimental, Mechan. Prop., 7)

[2001Bra] Braun, J., Ellner, M., “Phase Equilibria Investigations on the Aluminium-Rich Part of the

Binary System Ti-Al” Metall. Mater. Trans. A, 32A, 1037-1048 (2001) (Crys. Structure,

Equi. Diagram, Experimental, *, #, 34)

[2001Sun] Sun, F.-S., Kim, S.-E., Cao, Ch.-X., Lee, Y.-T., Yan, M.-G., “A Study of Ti5Si3/ Interface

in TiAl Alloys”, Scr. Mater., 45(4), 383-389 (2001) (Crys. Structure, Equi. Diagram,

Experimental, 11)

[2002Aze] Azevedo, C.R.F., Flower, H.M., “Experimental and Calculated Ti-Rich Corner of the

Al-Si-Ti Ternary Diagram”, Calphad, 26(3), 353-373 (2002) (Calculation, Crys. Structure,

Equi. Diagram, Experimental, Thermodyn., *, #, 52)

[2002Hok] Hokamoto, K., Lee, J. S., Fujita, M., Iton, S., Raghukandan, K., “The Synthesis Bulk

Material Through Explosive Compaction for Making Intermetallic Compound Ti5Si3 and

its Composites”, J. Mater. Sci., 37(19), 4073-4078 (2002) (Crys. Structure,

Experimental, 17)

8


MSIT®

Al–Si–Ti

[2002Sah] Saheb, N., Laoui, T., Daud, A.R., Yahaya, R., Radiman, S., “Microstructure and Hardness

Behaviours of Ti-Containing Al-Si Alloys”, Philos. Mag. A, 82(4), 803-814 (2002) (Crys.

Structure, Experimental, Mechan. Prop., 21)

[2002Sun] Sun, F.-S., Sam Froes, F.H., “Precipitation of Ti5Si3 Phase in TiAl Alloys”, Mater. Sci.

Eng. A, 328, 113-121 (2002) (Mechan. Prop., Experimental, 34)

[2003Luk] Lukas, H.-L., “Al-Si (Aluminium-Silicon)”, MSIT Ternary Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH,

Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 29)

[2003Sch] Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Ternary Evaluation Program, in

MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services,

GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 85)

Table 1: Crystallographic Data of Solid Phases

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Ti)(h)

1670 - 882

cI2

Im3m

W

a = 330.65 pure Ti [Mas2]

0 to 44.8 at.% Al at 1490°C [2003Sch]

0 to 4.7 at.% Si at 1330°C [1987Mur]

Possible ordering from A2 to B2 ( 2Ti)

[2003Sch]

( Ti)(r)

< 1490

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

pure Ti at 25°C [Mas2]

0 to 51.4 at.% Al [2003Sch]

0 to 0.5 at.% Si at 865°C [1987Mur]

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 pure Al at 25°C [Mas2]

0 to 0.6 at.% Ti [2003Sch]

(Si)

< 1414

cF8

Fd3m

C (diamond)

a = 543.06 at 25°C [Mas2]

Ti(Al1-xSix)3

TiAl3 (h)

1393 - 735

tI8

I4/mmm

TiAl3 (h) a = 384.9

c = 860.9

a = 378

c = 853.8

74.2-75.0 at.% Al in Al-Ti [2003Sch]

0 x 0.15.

D022 ordered phase at x = 0 [2001Bra]

at x = 0.15 [1965Ram]

TiAl3 (r)

< 950

tI32

I4/mmm

TiAl3 (r)

a = 387.7

c = 3382.8

74.5-75.0 at.% Al [2003Sch]

9


MSIT®

Al–Si–Ti

Ti2Al51416 - 990

Tetragonal

superstructure of

AuCu-type

[2001Bra]

tP28

P4/mmm

“Ti2Al5”

a = 395.3

c = 410.4

a = 391.8

c = 415.4

a = 390.53

c = 2919.63

chosen stoichiometry summarizing

several phases [2003Sch]:

Ti5Al11

66-71 at.% Al at 1300°C. Stable range

1416-995°C [2003Sch]

(including the stoichiometry Ti2Al5)

at 66 at.% Al [2003Sch]

at 71 at.% Al [2003Sch]

“Ti2Al5”

~1215-985°C; included in homogeneity

region of Ti5Al11 [2003Sch]

TiAl2< 1199

oC12

Cmmm

ZrGa2

tP4

P4/mmm

AuCu

tI24

I41/amd

HfGa2

tP32

P4/mbm

Ti3Al5

a = 1208.84

b = 394.61

c = 402.95

a = 403.0

c = 395.5

a = 397.0

c = 2497.0

a = 1129.3

c = 403.8

chosen stoichiometry summarizing

several phases [2003Sch]:

metastable modification of TiAl2, only

observed in as-cast alloys [2001Bra];

Ti1-xAl1+x; 63 to 65 at.% Al at 1250°C,

stable range 1445-1170°C

at 1300°C [2001Bra]

stable structure of TiAl2 < 1216

[2001Bra];

Ti3Al5, stable below 810°C

[2001Bra]

Ti3Al5< 810

tP32

P4/mbm

Ti3Al5

a = 1129.3

c = 403.8

,TiAl superstructure observed toward

37 at.% Ti [2003Sch]

Ti1-xAl1+x

1445 - 1170

oP4 a = 402.62

b = 396.17

c = 402.62

Probably metastable

0.26 < x < 0.31

at x = 0.28 [1990Sch]

, TiAl

< 1463

tP4

P4/mmm

AuCu(I)

a = 400.0

c = 407.5

a = 400.0

c = 407.5

a = 398.4

c = 406.0

46.7-66.5 at.% Al [2003Sch]

33.5 to 53.3 at.% Ti [1992Kat]

38 to 50 at.% Ti at 1200°C [2001Bra]

Ordered L10 phase

at 48.0 at.% Ti, [2001Sun]

at 50.0 at.% Al, [2003Sch]

at 62.0 at.% Al [2003Sch]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

10


MSIT®

Al–Si–Ti

2, Ti3Al

< 1164

(up to 10 GPa at RT)

hP8

P63/mmc

Ni3Sn a = 580.6

c = 465.5

a = 574.6

c = 462.4

~20 to 38.2 at.% Al [2003Sch]

D019 ordered phase

at 22 at.% Al [2003Sch]

38 at.% Al [2003Sch]

Ti3Si

< 1170

tP32

P42/n

Ti3P

a = 1039.0

c = 517.0

[1987Mur]

Ti5Si3< 2130

hP16

P63/mcm

Mn5Si3

a = 744.5

c = 514.6

a = 752.9

c = 525.0

a = 762.8

c = 527.1

35.5 to 39.5 at.% Si at 1920°C

[1987Mur, 1996Sei].

Dissolves up to 9 at.% Al [1997Bul]

Ti5(Si,Al)3 in equilibrium with ,TiAl

[2001Sun]

Ti65Al20Si15, metastable, obtained by

ball milling [1997Gua]

Ti5Si4< 1920

tP36

P41212

Zr5Si4

a = 713.3

c = 1297.7

[1987Mur]

TiSi

< 1570

oP8

Pnma

FeB

a = 654.4

b = 363.8

c = 499.7

[1961Bru]

TiSi2< 1500

oF24

Fddd

TiSi2

a = 825.3

b = 478.3

c = 854.0

[1987Mur]

* 1,

(Ti1-xAlx)8(AlySi1-y)16

Ti7Al5Si12

tI24

I41/amd

Zr3Al4Si5

a = 357.6 to 364.5

c = 2715 to 2865

a = 357

c = 2715

x 0.12 [1963Sch, 1965Ram]

0.06 y 0.25 at x = 0.12, y = 0.25

* 2,

Ti(AlxSi1-x)2

oC12

Cmcm

ZrSi2

a = c = 359.0 to

361.8

b = 1351.7

a = c = 360

b = 1353

0.15 x 0.3, pseudotetragonal

[1961Bru] at Ti31Al19Si50 (x 0.28)

[1965Ram]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

11


MSIT®

Al–Si–Ti

Table 2: Invariant Equilibria

Reaction T [°C] Type Phase Composition (at.%)

Al Si Ti

L ( Ti) + Ti5Si3 1545 e1 L

( Ti)

Ti5Si3

27

33

7

8

1

30

65

66

63

L + ( Ti) ( Ti) + Ti5Si3 1420 U1 L

( Ti)

( Ti)

Ti5Si3

47

45

48

9

4

0.9

0.5

30

49

54.1

51.5

61

L + ( Ti) + Ti5Si3 1415 U2 L

( Ti)

Ti5Si3

48

52

48

9

4

0.5

0.8

30

48

47.5

51.2

61

( Ti) Ti5Si3 + 2 + ~1035 E1 - - - -

( Ti) + Ti5Si3 ( Ti) + Ti3Si ~930 U3 - - - -

L + 1 (Si) + (Al) 579 U4 L

1

(Si)

(Al)

87.835

20.8

0

~98.5

12.1

50

100

~1.5

0.065

29.2

0

0

9

700°C

800°C

590°C

580°C (Si)(Al)

τ1

10 11 12 13 14

0.0

0.1

0.2

0.3

0.4

Si, at.%

Ti,

at.%

600°C

U4

Fig. 1: Al-Si-Ti.

Liquidus surface near

the Al-Si binary

eutectic

12


MSIT®

Al–Si–Ti

40

60

80

20 40 60

20

40

60

Ti Ti 30.00

Al 70.00

Si 0.00

Ti 30.00

Al 0.00

Si 70.00Data / Grid: at.%

Axes: at.%

Ti5(Si,Al)

3

Ti5(Si,Al)

3

(βTi)e

1

U1 U

2

(αTi)γ

e2

p2

p1(βTi)

γ

Fig. 2: Al-Si-Ti.

Liquidus and solidus

surfaces in the Ti-rich

corner. Solidus is

shown by dashed

lines

Fig. 3: Al-Si-Ti. Partial reaction scheme

Al-Ti A-B-C

l + (βTi) (αTi)

1490 p1 L (βTi)+Ti

5Si

3

1545 e1

Al-Si-Ti

L+(βTi) (αTi)+Ti5Si

31420 U

1

Si-Ti

l (βTi)+Ti5Si

3

1330 e2

l + (αTi) γ1462 p

2

(αTi) α2 + γ

1118 e3 (βTi)+Ti

5Si

3Ti

3Si

1170 p3

(βTi) (αTi)+Ti3Si

835 e4

L + (αTi) γ + Ti5Si

31415 U

2

(αTi) + Τi5Si

3α

2+ γ1035 U

3

(βTi)+Τi5Si

3(αTi)+Ti

3Si930 U

4

(αTi)+γ+Ti5Si

3

(αTi)+(βTi)+Ti5Si

3

(αTi)+Ti3Si+Ti

5Si

3

(αTi)+α2+Ti

5Si

3α

2+γ+Ti

5Si

3

L (αTi)+Ti5Si

3

?

13


MSIT®

Al–Si–Ti

40

60

80

20 40 60

20

40

60

Ti Ti 30.00

Al 70.00

Si 0.00

Ti 30.00

Al 0.00


Axes: at.%

L+Ti5(Si,Al)

3

Ti5(Si,Al)

3

L+(βTi)+Ti5(Si,Al)

3

(βTi)+Ti5(Si,Al)

3

L+(βTi)+Ti5(Si,Al)

3

L+Ti5(Si,Al)

3

L+(βTi)

L

L

L+(βTi)

(βTi)

Fig. 4: Al-Si-Ti.

Calculated partial

isothermal section at

1523°C

40

60

80

20 40 60

20

40

60

Ti Ti 30.00

Al 70.00

Si 0.00

Ti 30.00

Al 0.00


Axes: at.%

γ

Ti5(Si,Al)

3

(αTi)+(βTi)+Ti5(Si,Al)

3

(αTi)+Ti5(Si,Al)

3

(αTi)+γ+Ti5(Si,Al)

3

(βTi)

(βTi)+Ti5(Si,Al)

3

(αTi)

Fig. 5: Al-Si-Ti.

Isothermal section at

1300°C [1997Bul]

14


MSIT®

Al–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Al

Si Data / Grid: at.%

Axes: at.%

TiSi2

TiSi

Ti5Si

4

Ti5Si

3

(βTi) (αTi) γ Ti2Al

5 TiAl3(h)

L

τ2

(Si)Fig. 6: Al-Si-Ti.


1200°C

20

40

60

80

20 40 60 80

20

40

60

80

Ti Al


Axes: at.%

TiSi2

TiSi

Ti5Si

4

Ti5Si

3

Ti3Si

(αTi) α2 TiAl TiAl

2 TiAl3(r)

L

τ1

τ2

possible range of τ3

(Si)Fig. 7: Al-Si-Ti.


700°C

15


MSIT®

Al–Si–Ti

1000°C

1200°C

800°C

0

10 20 30 40

1

2

3

4

5

Ti

Al, at.%

Si,

at.

%

Ti Si3 Ti Si3Ti Si5 3

Fig. 8: Al-Si-Ti.

Isothermal sections of

the Ti-rich corner at

1200, 1000 and

partially at 800°C

16


MSIT®

Al–Sn–Ti

Aluminium – Tin – Titanium

Anatoliy Bondar, Olga Fabrichnaya

Literature Data

The Ti rich corner of the system has been of great interest owing to the fact that Ti alloys with additions of

Al and Sn are widely used in industry, in particular, the alloy Ti-5Al-2.5Sn (mass%, Ti to balance)

[1954Fin, 1957Jaf, 1990Lam, 2003Mat]. Properties of the Ti-5Al-2.5Sn alloy have been the subject of

many articles in the literature. The known phase equilibria of the Ti rich corner are due mainly to the work

of [1960Kor, 1961Kor, 1962Kor] (isothermal sections at 600, 1000, and 1200°C for the region

Ti-Ti3Sn-Ti45Al55) and [1969Cro] (alloys annealed at 600-700°C containing (8-13)Al-(1-2)Sn (at.%)). The

data have been assessed in [1966Gla] and [1993Kub]. [1993Pie, 1997Pie] have presented an isothermal

section for the entire system at 900°C.

[1960Kor, 1961Kor, 1962Kor, 1963Kor] used both iodide and “magnesiumthermic” Ti, Al of 99.99 %

purity and Sn of 99.9 % purity as their starting materials. The majority of their ingots were melted in a

nonconsumable electrode arc furnace and the remainder using levitation induction melting. Alloys

containing up to 22 mass% Al+Sn were forged and homogenized at 1200°C under vacuum for 100 h. The

samples were then water quenched after heating at 1200°C for 75 h, or at 1000°C for 200 h. Also, samples

were annealed at 1100°C for 50 h, at 1000°C for 200 h, at 800°C for 300 h and at 600°C for 500 h followed

by furnace cooling. Other samples were annealed at 1000°C for 500 h, at 800°C for 1000 h and at 600°C

for 1200 h followed by furnace cooling. The alloys were studied using light microscopy, XRD (in some

cases) and thermal analysis using Nedumov's apparatus [1960Ned] (using the temperature dependence of

sample conductance). In their studies, Kornilov & Nartova did not distinguish between the ( Ti) phase and

the Ti3(Al,Sn) solid solution, the XRD patterns of which differ only by weak superstructure lines appearing

in the pattern of the latter.

[1969Cro] prepared alloys from high purity electrolytically refined Ti and alloying additions of at least

99.9 % purity in a nonconsumable electrode arc furnace. The alloys were hot rolled at 1000°C and then

annealed at 900°C for 48 h, at 800°C for 200 h, at 700°C for 500 h and at 600°C for 1000 h followed by

quenching in iced bulbs. The microstructures of the deeply etched samples (swabbing with 0.5 % HF, 1.5 %

HNO3 in a saturated aqueous solution of citric acid followed by only the saturated aqueous solution of citric

acid) were examined to reveal the presence of particles of the 2 phase.

[1984Li] prepared alloys from Ti of 99.9 mass% purity, Al of 99.999 % and Sn of 99.9 %. The samples were

annealed at 600°C for 400 h and studied by optical microscopy.

[1993Pie] prepared alloys by arc melting powdered elements (Ti of 99 % purity, Al of 99.8 % and Sn of

99.5 %). The samples were annealed at 900°C for 140 h and water quenched before examination by powder

XRD using Guinier-Huber cameras and Cu K 1 radiation.

[1994Kus] studied alloys of Ti-(50-52)Al-(0-5)Sn (at.%) in the as cast and annealed (at 1000°C for 168 h)

conditions by optical and electron transmission microscopy, XRD and Vickers hardness measurement. The

alloys contained 0.3-0.4 % O, 0.1-0.2 % N, 0.04-0.05 % C, 0.04-0.07 H, and 0.04-0.06 % Fe (at.%).

Binary Systems

The Sn-Ti and Al-Sn binary systems are accepted from [Mas2] (where the Sn-Ti phase diagram was taken

from [1987Mur]). For the Al-Ti system, critical assessments of [2003Sch] and [2003Gry] are available,

which are based on the latest thermodynamic optimizations of [1992Kat, 1997Zha] and [2000Ohn] and

taking into account experimental studies. The Al-Ti phase diagram was accepted from [2003Gry] which

combined the liquidus and the Ti rich part (up to 50 at.% Al) of the diagram presented by [1997Zha], the

work of [2000Ohn] (the CsCl type ordered region within the ( Ti) field) and also solid phase equilibria in

the range 50 to 75 at.% Al as presented by [2001Bra].

17


MSIT®

Al–Sn–Ti

Solid Phases

The binary phases, relevant to the phase equilibria under consideration and the ternary Ti5Al2Sn phase are

listed in Table 1. The only ternary phase reported is given by [1993Pie] and has a narrow homogeneity

range, based on the observed variation in the lattice parameters.

The isostructural 2,Ti3Al and Ti3Sn phases have been shown by [1961Kor, 1962Kor] to form a complete

series of solid solutions Ti3(Al,Sn). As found in [1993Pie], Ti5Sn3 dissolves up to 18 at.% Al and Ti6Sn5

dissolves up to 3 at.% Al. There are conflicting opinions on the Sn solubility in ,TiAl; up to 18 mass% Sn

at 22 mass% Al(Ti56Al37Sn7) in [1960Kor, 1961Kor, 1962Kor] at 600, 1000 and 1200°C, up to the

composition Ti45Al51Sn4 at 1000°C in [1994Kus] and no solubility at all given by [1993Pie].

The site occupancies of Sn in Ti3Al in an alloy of Ti-26Al-(1-2)Sn, and in ,TiAl in an alloy of Ti-51Al-3Sn

(at.%) were measured by the atom location channeling enhanced microanalysis (ALCHEMI) method by

[1999Hao]. In the both phases, Sn atoms were found to occupy Al sites. However, [1994Kus] reported a

more complicated influence of Sn alloying on site occupation in the phase. Tin atoms occupy both Ti

(predominantly) and Al sites, and with increasing Sn content, the mutual exchanges of Ti and Al atoms

increase in the lattice sites. The solubility lobes of Sn in ,TiAl, as seen in the isothermal sections given by

Kornilov & Nartova [1960Kor, 1962Kor] (Figs. 2, 4, 5), showing some extension of the homogeneity range

towards increasing Ti content, are in agreement with the data of [1994Kus].

Liquidus and Solidus Surfaces

The liquidus and solidus surfaces of the Al-Sn-Ti system in the Ti-rich corner have been estimated by

[2000Bul] on the basis of experimental data [1961Kor, 1962Kor, 1969Cro] and the associated binary

systems. They are shown in Fig. 1. The ( Ti)+ 2 eutectic was observed microstructurally by [1962Kor] to

lie at 45 mass% Ti3Sn (Ti3Al to balance, i.e. at Ti75Al16Sn9), the melting point being a little lower than the

binary eutectic le ( Ti)+Ti3Sn (1605°C after [1987Mur]).

Isothermal Sections

Isothermal sections at 600, 1000 and 1200°C were constructed in [1960Kor, 1962Kor] for the Ti-rich

portion of the system. In the 600°C isothermal section the narrow two phase field of ( Ti)+ 2 was found to

adjoin Sn-Ti side. This result of [1960Kor, 1962Kor] contradicts later data of [1969Cro] and [1984Li], as

well as the Sn-Ti and Al-Ti binary phase diagrams. The isothermal section at 600°C was modified by

[1993Kub] taking into account data of [1969Cro] and [1984Li] for the ( Ti)/( Ti)+ 2 phase boundary and

phase diagrams of the binary systems. The isothermal section at 600°C is presented in Fig. 2.

The 900°C isothermal section presented by [1993Pie, 1997Pie] is shown in Fig. 3 with some modifications

taking into account the solubility of Sn in ,TiAl according to the work of [1962Kor, 1994Kus] and the

binary systems. The authors of [1993Pie] noted an absence of Sn solubility in the ,TiAl phase but they did

not give any composition. The alloys near the Al-Sn side are in the liquid state at 900°C and the liquid is in

equilibrium with the Ti6Sn5 and TiAl3 solid phases. The ternary phase Ti5Sn2Al has been found by

[1993Pie] at 900°C, but the temperature range of its stability is not reported. If its stability range is wide

enough, the appearance of the ternary phase could influence phase equilibria in the isothermal sections

presented in Figs. 2, 4, 5.

The isothermal sections at 1000 and 1200°C are shown according to [1962Kor] taking into account data for

the binary systems (Figs. 4, 5). The homogeneity range of the ( Ti) phase at 1000 and 1200°C is delineated

by dashed lines. It should be noted that the data of [1994Kus] for the solubility of Sn in ,TiAl as

Ti56Al37Sn7 agrees with the 1000°C isothermal section (Fig. 4). The Sn influence on the ( Ti)/( Ti)+ 2

phase boundary was studied by Crossley [1969Cro], also at 700°C and Sn contents up to 2 at.%, where the

phase boundary was found to be at ~13 at.% Al.


[1961Kor, 1962Kor] presented the section through the Ti corner at an equal ratio of Al to Sn in mass%

(Al:Sn=1:1) and the section Ti3Al-Ti3Sn (Figs. 6, 7, respectively). There is a remarkable difference in the

18


MSIT®

Al–Sn–Ti

composition for the / + 2 boundary presented in the isothermal sections and the isopleth. The Ti corner

at Al:Sn=1:1 is corrected to be in agreement with the isothermal sections. The portion of the Ti3Al-Ti3Sn

section at temperatures below 1200°C is shown by dashed lines because the equilibria involving ( Ti) were

omitted by [1961Kor, 1962Kor] as already discussed above.

Thermodynamics

The heat capacity of the Ti-5Al-2.5Sn alloy is reported for temperatures between 4 and 290 K [1978Ili] and

between 273 and 973 K [1986Ric].


The commercial alloy Ti-5Al-2.5Sn is widely used and there are many references to its properties in the

literature [1988Fuj, 1990Lam, 2003Mat], including low temperature properties [1980Kaw, 1993Gri1,

1993Gri2, 2001Sun] and cyclic loading behavior [2001Sun].

[1954Fin] showed that Sn additions of up to 5 mass% to the (0-5)Al-Ti (in mass%) alloys led to an increase

in bend and tensile strength properties without suffering any loss in hot fabricability or substantial loss of

ductility.

[1963Kor] reported results of bending creep tests for alloys in the section of equal Al to Sn ratio (in mass%

Al:Sn=1:1) and sections Ti3Al-Ti3Sn and TiAl-Ti3Sn. Other properties were also studied. The creep

behavior was studied at 700°C using a technique presented in [1957Pro]. The maximum creep resistance

was found in the alloy of composition Ti3Al:Ti3Sn=1:1 (in mass%) and alloys based on the phase.

Miscellaneous

In the literature, there is information concerning the hydrogen solubility of the alloy Ti-5Al-2.5Sn

[1958Alb] and the influence of hydrogen on its properties [1972Wil, 1984Ham].

References

[1954Fin] Finlay, W.L., Jaffee, R.I., Parcel, R.W., Durstein, R.C., “Tin Increases Strength of Ti-Al

Alloys without Loss in Fabricability”, J. Metals, 6(1), 25-29 (1954) (Mechan. Prop.,

Experimental, 5)

[1957Jaf] Jaffee, R.I., Ogden, H.R., Maykuth, D.J., “Al-Sn-Ti Alloys with , and Compound

Formers”, Pat. USA 2779677, (1957) (Mechan. Prop., Experimental)

[1957Pro] Prokhanov, V.F., “New Model of Machine for High-Temperature Strength Test with

Centrifugal Technique” (in Russian), Zavods. Lab., 23(8), 983-984 (1957) (Mechan. Prop.,

Experimental, 1)

[1958Alb] Albrecht, W.M., Mallett, M.W., “Hydrogen Solubility and Removal for Titanium and

Titanium Alloys”, Trans. Met. Soc. AIME, 212, 204-210 (1958) (Corrosion,

Experimental, 10)

[1960Kor] Kornilov, I.I., Nartova, T.T., “Phase Diagram of the Ti-Al-Sn System” (in Russian), Dokl.

Akad. Nauk SSSR, 131(4), 837-839 (1960) (Equi. Diagram, Experimental, 8)

[1960Ned] Nedumov, N.A., “High-Temperature Method of Contact-less Thermography” (in Russian),

Zh. Phiz. Khim., 34(1), 184-191 (1960) (Phys. Prop., Experimental, 13)

[1961Kor] Kornilov, I.I., Nartova, T.T., “Continuous Solid Solutions of the Metallides Ti3Al-Ti3Sn in

the Ti-Al-Sn System” (in Russian), Dokl. Akad. Nauk SSSR, 140(4), 829-831 (1961) (Equi.

Diagram, Electr. Prop., Mechan. Prop., Experimental, 16)

[1962Kor] Kornilov, I.I., Nartova, T.T., “Phase Diagram of the Titanium-Aluminium-Tin System”, in

“Titanium and its Alloys. Issue 7. Metallochemistry and New Alloys” (in Russian), Ageev,

N.B., Kornilov, I.I., Fedotov, S.G. (Eds.), Akad. Nauk SSSR, Moscow, 95-104 (1962)

(Equi. Diagram, Experimental, 16)

[1963Kor] Kornilov, I.I., Nartova, T.T., “Investigation of High-Temperature Strength of the Ti-Al-Sn

Alloys”, in “Titanium and its Alloys. Issue 10. Investigation of Titanium Alloys” (in

19


MSIT®

Al–Sn–Ti

Russian), Kornilov, I.I., Pylaeva, E.N., Boriskina N.G. (Eds.), Akad. Nauk SSSR, Moscow,

202-206 (1963) (Mechan. Prop., 6)

[1966Gla] Glazova, V.V., “Alloying of Titanium” (in Russian), Metallurgiya, Moscow, (1966) 170-176

(Equi. Diagram, Mechan. Prop., Review, 293)

[1969Cro] Crossley, F.A., “Effects of the Ternary Additions: O, Sn, Zr, Cb, Mo, and V on the

/ +Ti3Al Boundary of Ti-Al Base Alloys”, Trans. Metall. Soc. AIME, 245(9), 1963-1968

(1969) (Equi. Diagram, Experimental, 15)

[1972Wil] Wiliams, D.P., Nelson, H.G., “Gaseous Hydrogen-Induced Cracking of Ti-5Al-2.5Sn”,

Met. Trans. (J. of Metals, AIME), 3(8), 2107-2113 (1972) (Corrosion, Experimental, 29)

[1978Ili] Iliyev, L.B., Ovcharenko, V.I., Pervakov, V.A., “Low-Temperature Heat Capacity of

Commercial Grade Titanium VT1-0 and Its Alloys VT5 and VT5-1”, Phys. Met. Metall.

(Engl. Transl.), 46(4), 34-39 (1978), translated from Fiz. Met. Metallaved., 46(4), 719-725

(1978) (Thermodyn., Experimental, 13)

[1980Kaw] Kawabata, T., Morita, Sh., Izumi, O., “Deformation and Fracture of Ti-5Al-2.5Sn ELI Alloy

at 4.2K~291K”, in Titanium 80: Sci. Technol. Proc. 4 Int. Conf., 2, 801-809 (1980)

(Mechan. Prop., Experimental, 4)

[1984Ham] Hammond, C., Spurling, R.A., Paton, N.E., “Hydride Precipitation and Dislocation

Substructures in Ti-5 Pct Al-2.5 Pct Sn”, Metall. Trans. A, A15(1-6), 813-817 (1984)

(Corrosion, Experimental, 7)

[1984Li] Li, D., Liu, Y., Wan, X., “On The Thermal Stability of Titanium Alloys. I. The Electron

Concentration Rule for Formation of Ti3X Phase” (in Chinese), Acta Metall. Sin. (Jinshu

Xuebao), 20(6), A375-A383 (1984) (Equi. Diagram, Crys. Structure, Experimental,

Calculation, 22)

[1986Ric] Richter, F., Born, L., “Specific Heat Capacities of Metallic Materials, Part III: Five

Non-Ferrous Metals, Including NiCr15Fe (INCONEL 600)” (in German), Z. Werkstofftech.,

17(7), 233-237 (1986) (Thermodyn., Experimental, 12)

[1987Mur] Murray, J.L., “The Sn-Ti (Tin-Titanium) System”, in “Phase Diagrams of Binary Titanium

Alloys”, Murray, J.L., (Ed.), ASM, Metals Park, Ohio (1987) 294-299 (Equi. Diagram, Crys.

Structure, Thermodyn., Assessment, 22)

[1988Fuj] Fujii, H., “Characteristics of Continuous Cooling Transformation in Ti-5Al-2.5Sn” (in

Japanese), Curr. Adv. Mater. Process, 1(2), 399 (1988) (Mechan. Prop., Experimental, 2)

[1990Lam] Lampman, S., “Wrought Titanium and Titanium Alloys”, Metals Handbook, Tenth Edition.

Vol. 2. Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, 2,

592-633 (1990) (Mechan. Prop., Review, 32)


Ti-Al System”, Metall. Trans. A, 23A (8), 2081-2090 (1992) (Equi. Diagram, Thermodyn.,

Assessment, Calculation, 51)

[1993Gri1] Grinberg, N.M., Aleksenko, E.N., Moskalenko, V.A., Smirnov, A.R., Yakovenko, L.F.,

Mozhaev, A.V., Arinushkin, I.A., “Fatigue-Induced Dislocation Structure of Titanium

Alloy VT5-1ct at Temperatures of 293-11 K”, Mater. Sci. Eng. A, A165(2), 117-124 (1993)

(Crys. Structure, Mechan. Prop., Experimental, 14)

[1993Gri2] Grinberg, N.M., Smirnov, A.R., Moskalenko, V.A., Aleksenko, E.N., Yakovenko,

Zmievsky, V.I., “Dislocation Structure and Fatigue Crack Growth in Titanium Alloy

VT5-1ct at Temperatures of 293-11 K”, Mater. Sci. Eng. A, A165(2), 125-131 (1993) (Crys.

Structure, Mechan. Prop., Experimental, 35)

[1993Kub] Kubaschewski, O., “Aluminium - Tin - Titanium”, MSIT Ternary Evaluation Program, in

MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services

GmbH, Stuttgart; Document ID: 10.15979.1.20 (1993) (Equi. Diagram, Crys. Structure,

Assessment, 6)

[1993Pie] Pietzka, M.A., Gruber, U., Schuster, J.C., “Investigation of Phase Equilibria in the Ternary

Ti-Al-Sn”, J. Phys. IV, Coll. C7, Suppl. J. Phys. III, 3, 473-476 (1993) (Equi. Diagram,

Crys. Structure, Experimental, 2)

20


MSIT®

Al–Sn–Ti

[1994Bra] Braun, J., Ellner, M., Predel, B., “On the Structure of the High Temperature Phase

Ti1-xAl1+x” (in German), J. Alloys Compd., 203, 189-193 (1994) (Equi. Diagram, Crys.

Structure, Experimental, 8)

[1994Kus] Kusabiraki, K., Yamamoto, Y., Ooka, T., “Effects of Tin Addition on Microstructure and

Crystal-Structure of TiAl-Base Alloys” (in Japanese), Tetsu To Hagane - J. Iron Steel Inst.

Jpn., 80(10), 67-71 (1994) (Crys. Structure, Mechan. Prop., Experimental)

[1995Pie] Pietzka, M.A., Schuster, J.C., “New Ternary Aluminides T5M2Al Having W5Si3-Type

Structure”, J. Alloys Comp., 230, L10-L12 (1995) (Crys. Structure, Experimental, 7)

[1997Pie] Pietzka, M.A., Schuster, J.C., “Phase Equilibria of the Quaternary System Ti-Al-Sn-N at

900°C”, J. Alloys Comp., 247, 198-201 (1997) (Equi. Diagram, Crys. Structure,

Experimental, 12)

[1997Zha] Zhang, F., Chen, S.L., Chang, Y.A., Kattner, U.R., “A Thermodynamic Description of the

Ti-Al System”, Intermetallics, 5, 471-482 (1997) (Equi. Diagram, Thermodyn.,

Experimental, Assessment, 45)

[1999Hao] Hao, Y.L., Xu, D.S., Cui, Y.Y., Yang, R., Li, D., “The Site Occupancies of Alloying

Elements in TiAl and Ti3Al Alloys”, Acta Mater., 47(4), 1129-1139 (1999) (Crys.


[2000Bul] Bulanova, M., Ban’kovsky, O., Soroka, A., Samelyuk, A., Tretyachenko, L., Kulak, L.,

Firstov, S., “Phase Composition, Structure and Properties of Cast Ti-Si-Sn-Al Alloys”,

Z. Metallkd., 91, 64-70 (2000) (Crys. Structure, Experimental, Mechan. Prop., Equi.

Diagram, 22)

[2000Ohn] Ohnuma, I., Fujita, Y., Mitsui, H., Ishikawa, K., Kainuma, R., Ishida, K., “Phase Equilibria

in the Ti-Al Binary System”, Acta Mater., 48, 3113-3123 (2000) (Equi. Diagram, Crys.

Structure, Thermodyn., Experimental, Assessment, 37)

[2001Sun] Sun, Q.Y., Gu, H.C., “Tensile and Low-Cycle Fatigue Behaviour of Commercially Pure

Titanium and Ti-5Al-2.5Sn Alloy at 293 and 77 K”, Mater. Sci. Eng. A, A316, 80-86 (2001)

(Mechan. Prop., Experimental, 12)


Binary System Ti-Al”, Metall. Mater. Trans. A, 32A, 1037-1048 (2001) (Equi. Diagram,

Crys. Structure, Experimental, Review, 34)

[2003Gry] Grytsiv, A., Rogl, P,. Schmidt, H., Giester, G., “Constitution of the Ternary System Al-Ru-

Ti (Aluminium - Ruthenium - Titanium)”, J. Phase Equilib., 24(6), 511-527 (2003) (Equi.

Diagram, Crys. Structure, Experimental, Review, 45)

[2003Mat] “MatWeb: Materials Property Data”, http://www.matweb.com/ (Mechan. Prop.,

Review, 2)

[2003Sch] Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Service, GmbH,

Stuttgart, to be published, (2003) (Equi. Diagram, Review, 85)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96 at 25°C [Mas2]

(Sn)

< 231.9681

tI4

I41/amd

Sn

a = 583.15

c = 318.14

[Mas2]

21


MSIT®

Al–Sn–Ti

( Ti)

< 1490

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

[Mas2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 [Mas2]

2,Ti3SnxAl1-x

< 1670

hP8

P63/mmc

Ni3Sn a = 591.6

c = 476.4

a = 577.5

c = 465.5

0 < x < 1

Ti3Al is stable at T < 1166

x = 1 [1987Mur]

x = 0 [V-C2]

Ti2Sn

< 1550

hP6

P63/mmc

InNi2

a = 465.3

c = 570

[1987Mur]

Ti5Sn3

1510

hP16

P63/mcm

Mn5Si3

a = 804.9

c = 545.4

[1987Mur]

Ti6Sn5

~1490 - 790

hP22

P63/mmc

Ti6Sn5

hP22

P31c

?

a = 922

c = 569

a = 924.8

c = 589

[1987Mur]

[1987Mur]

Ti6Sn5

< 790

oI44

Immm

Nb6Sn5

a = 1693

b = 914.4

c = 573.5

[1987Mur]

,TiAl

< 1463

tP4

P4/mmm

AuCu

a = 400.0

c = 407.5

a = 398.4

c = 406.0

a = 399.6

c = 407

a = 399.5

c = 408

a = 400

c = 408.6

a = 399.8

c = 408.9

at 50.0 at.% Al [2001Bra]

at 62.0 at.% Al [2001Bra]

at 50.55 at.% Al [1994Kus]

at 51.89 at.% Al [1994Kus]

at 50.56Al-2.25Sn (at.%) [1994Kus]

at 51.12Al-3.92Sn (at.%) [1994Kus]

TiAl2< 1215

tI24

I41/amd

HfGa2

a = 0.3970

c = 2.4970

[2001Bra]

TiAl3 (h)

1387 - 735

tI8

I4/mmm

TiAl3 (h)

a = 384.9

c = 860.9

[2001Bra]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

22


MSIT®

Al–Sn–Ti

TiAl3 (l)

< 950 (Ti rich)

tI32

I4/mmm

TiAl3 (l)

a = 387.7

c = 3382.8

[2001Bra]

TiAl3metastable

cP4

Pm3m

AuCu3

a = 397.2 0.1 [1994Bra]

* 1, Ti5Sn2Al

(obtained at 900°C)

tI32

I4/mcm

W5Si3

a = 1054.9 0.2

c = 524.2 0.2

a = 1056.4

c = 526.4

[1993Pie, 1995Pie]

Al rich composition;

Al lean

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

60

70

80

90

10 20 30 40

10

20

30

40

Ti Ti 50.00

Al 50.00

Sn 0.00

Ti 50.00

Al 0.00

Sn 50.00Data / Grid: at.%

Axes: at.%

(βTi)

α2

e1

Fig. 1: Al-Sn-Ti.

Partial liquidus and

solidus surfaces in Ti

corner

23


MSIT®

Al–Sn–Ti

50

60

70

80

90

10 20 30 40 50

10

20

30

40

50

Ti Ti 40.00

Al 60.00

Sn 0.00

Ti 40.00

Al 0.00


Axes: at.%

(αTi)

α2

γ

Fig. 2: Al-Sn-Ti.

Partial isothermal

section at 600°C

40

60

80

20 40 60

20

40

60

Ti Ti 20.00

Al 80.00

Sn 0.00

Ti 20.00

Al 0.00


Axes: at.%

Ti5Sn

2Al

(αTi)

Ti2Sn

βTi6 Sn

5 +TiAl3 +L

Ti5Sn

3

Ti3Al

γ

TiAl2

TiAl3

(βTi)

Ti3Sn

α2

βTi6Sn

5

Fig. 3: Al-Sn-Ti.

Partial isothermal

section at 900°C

24


MSIT®

Al–Sn–Ti

50

60

70

80

90

10 20 30 40 50

10

20

30

40

50

Ti Ti 40.00

Al 60.00

Sn 0.00

Ti 40.00

Al 0.00


Axes: at.%

(αTi)

α2

γ(βTi)

50

60

70

80

90

10 20 30 40 50

10

20

30

40

50

Ti Ti 40.00

Al 60.00

Sn 0.00

Ti 40.00

Al 0.00


Axes: at.%

(αTi) γ(βTi)

α2

Fig. 4: Al-Sn-Ti.

Partial isothermal

section at 1000°C

Fig. 5: Al-Sn-Ti.

Partial isothermal

section at 1200°C

25


MSIT®

Al–Sn–Ti

10 20

800

900

1000

1100

1200

1300

1400

1500

1600

1700

Ti 75.00

Al 25.00

Sn 0.00

Ti 75.00

Al 0.00

Sn 25.00Sn, at.%

Te

mp

era

ture

, °C

L

L+(βTi)

L+α2

(βTi) α2

(βTi)+α2

L+(βTi)+α2

1670°C

(αTi)

(αTi)+α2

(αTi)+α2+(βTi)

90 80 70 60

750

1000

1250

1500

1750

Ti Ti 57.94

Al 34.27

Sn 7.79Ti, at.%

Te

mp

era

ture

, °C

(αTi)

(βTi)+α2

(βTi)

L

L+(βTi)

882°C

1670°C

α2

Fig. 7: Al-Sn-Ti.

Vertical section

Ti3Al-Ti3Sn

[1961Kor, 1962Kor]

Fig. 6: Al-Sn-Ti.

Vertical section at the

equal Al to Sn ratio in

mass%

26


MSIT®

Al–Ti–V

Aluminium – Titanium – Vanadium

Ludmila Tretyachenko

Literature Data

Results of studies of phase equilibria in the Al-Ti-V system published approximately till 1991 were

reviewed by [1993Hay, 1995Hay] and used to present the ternary phase diagram as a number of isothermal

sections from the liquidus down to 800°C and partial isothermal sections of Ti rich corner of this phase

diagram from 900 down to 600°C. The reaction scheme was proposed from data on the structure of alloys

obtained by various authors while the experimental study of phase relations during crystallization of alloys

has not been carried out. It should be noted that the earlier data on the phase equilibria in the region

adjoining to the Al-Ti and Al-V sides were interpreted in accordance with versions of the Al-Ti and Al-V

phase diagrams, which differ from that accepted at present. So, the earlier experimental data were

reinterpreted taking into account new data on the Al-Ti system.

No ternary phases, continuous solid solutions between TiAl3 and VAl3 ( ), a significant solubility of V in

TiAl( ), a wide region of bcc solid solutions ( ), which undergo 2 (Ti3Al) transformations in a

region adjacent to the Al-Ti side, were reported. Four invariant equilibria involving a liquid were suggested

to exist. One of them, L+ + (V5Al8) at 1390°C, was shown from [1970Vol]. The phase fields + +

and + + were accepted to exist at high temperatures and transformed into + + and + + at a

temperature between 1100 and 1000°C.

The first attempt to determine the liquid - solid phase equilibria was made by [1991Par]. Thirty eight

samples prepared by induction melting were annealed at 900°C for 500 h and quenched in ice water. A study

of annealed and as cast alloys was made using optical microscopy and X-ray diffraction analysis (XRD).

Several alloys were examined by scanning electron microscopy (SEM) and electron microprobe analysis

(EMPA) using energy dispersive X-ray analysis (EDXA). The differential thermal analysis (DTA) of

several alloys was performed in the range of 1000 to 1400°C during heating and cooling at a rate of

25°C min-1 and sometimes during cooling with a furnace. Compositions of the phases coexisting in

three-phase alloys were used to determine the positions of the three-phase fields + 2+ , 2+ + , + +

and + + . Four invariant reactions were determined for the liquid-solid equilibria. Three of them were

found different from those proposed by [1993Hay, 1995Hay]. Moreover, the temperatures of the invariant

reactions in the Al-V system were accepted by [1991Par] from the earlier work of [1954Ros] rather than

from [1955Car] accepted in the assessment by [1989Mur] because the data from [1954Ros] were found to

be comparable with the melting temperature determined by [1991Par].

[1992Ahm, 1994Ahm1] studied 36 alloys prepared by arc melting, solution treated at 1200°C for 24 h,

quenched and then annealed at 900, 800, 700 and 600°C for 1.21, 2.42, 3.63 and 4.84 Ms respectively. The

samples were examined by means of optical microscopy, scanning and transmission electron microscopy

(SEM and TEM), XRD and EDXA. The isothermal sections at 1200, 900, 800, 700 and 600°C were

constructed. The ordering transformation bcc A2 ( ) B2 ( 0) was established to take place in a wide range

of concentration. The as cast alloys with 10 - 57 mass% Al and 4 -46 mass% V were studied by the same

method [1994Ahm2]. Cooling rate during solidification was estimated to be in the range of 10 to 100 K s-1.

The obtained experimental results were used to establish the liquidus projection of the Al-Ti-V system and

to analyze the solidification behavior of the alloys. Five invariant reactions were suggested to take place

during solidification. Two different phases, (Ti,V)Al3 and (V,Ti)Al3, instead of continuous solid solutions

were accepted to exist at temperatures 1200°C, but experimental data on crystal structure of these phases

were not reported.

A number of investigations were concerned to certain region of the ternary phase diagram.

[1995Ahm] studied aging behavior of the Ti-21.0V-39.5Al alloy (here and further compositions are given

in at.%, if not indicated differently). The alloy was arc melted, homogenized at 1200°C for 24 h and aged

at 600 and 700°C for various times from 3 to 96 h. The samples have been examined using calorimetric

differential thermal analysis (CDTA) at heating and cooling rates of 30, 20, 10 and 5 K min-1, optical

27


MSIT®

Al–Ti–V

microscopy, XRD, TEM and microhardness measurements. The transformation sequence on heating and

cooling was established in accordance with the calculated section for 21 at.% V.

The effect of V additions up to 7.5 mass% on the structure of Ti alloys with 34, 36, and 38 mass% Al in as

cast and annealed at 1000°C conditions was studied by means of optical microscopy, EMPA and XRD

[1988Has].

Alloying mechanism of V in TiAl based alloys were examined by [2001Sun]. The alloys 52Ti-xV-48Al

(x = 0 to 6) were prepared by arc melting under hot isostatic pressing at 1200°C for 3 h, heat treated at

1200°C for 12 h, air cooled at room temperature (RT) and aged at 800°C for 8 h. Then they have been

studied by optical microscopy, SEM, EMPA, TEM. Tensile tests at RT and 900°C and creep tests at 800°C

were carried out and the lattice parameters of the phase were determined.

In situ high temperature X-ray diffraction (HTXRD) was used by [1992Cha1] to study alloys in the 2+ +

phase field at 800, 1000 and 1100°C. [1992Cha2] has applied differential scanning calorimetry (DSC) and

HTXRD to examine phase transformations in the Ti-44Al-4, 7 and 15V alloys over the temperature range

20 - 1500°C.

The + equilibrium at 1300°C and 1200°C and 2+ at 1000°C have been studied by [2000Kai]. The

alloys Ti-(0.5-12)V-(35-47)Al were prepared by arc melting and heat treated at 1000°C and at 1200°C for

168 h, at 1300°C for 24 h. Chemical analysis has shown a high level of impurities in the as cast alloys with

V (2150 mass ppm O, 199 ppm N in the Ti-10V-47Al alloy), that was attributed to a high level of impurities

in the starting V material. The alloys were examined by optical microscopy and EMPA. The phase

boundaries of the ( 2)+ region at 1000, 1200 and 1300°C have been established.

Phase transformations in the arc melted Ti-10V-40Al (nominal composition) alloy, Ti-9.76V-41.73Al by

chemical analysis, have been studied by [1995Sha1], who used TEM, EMPA, SEM. The obtained

experimental results were discussed and compared with the calculated section of the Al-Ti-V system at 49

at.% Ti. The calculated liquidus projection and isothermal section at 900°C were presented. Details of the

calculations were not reported. [1997Sha] has summarized the results on the microstructure and the phase

transformations in the as cast Ti-10V-40Al alloy and has presented the calculated section at 50 at.% Ti.

The microstructure of the as cast Ti-50Al-20V alloy ( + ) was examined by [1995Sha3] using TEM, SEM,

EDXA. Ordering and decomposition of the phase in melt spun and aged ribbons of TiVxAl1-x (x = 0.2,

0.3, 0.5). Aging was performed at 400, 500 and 700°C for times up to 300 h. Ordering of phase, and 2

phases precipitation and metastable phase formation were considered. The phase transformations are

discussed from the calculated isopleth at 50 at.% Ti and the isothermal section at 900°C. The two-sublattice

descriptions were used for the disordered and the ordered B2 ( 0) phases to model the B2 ordering

transformation.

The calculated isothermal section at 600°C is presented by [2002Dip] to discuss an alloying behavior in the

Al-Ti-V system considering phenomena taking place on the atomic and subatomic level from information

obtained by electron spectroscopy.

An effect of V on the structure and properties of TiAl ( ) based alloys were studied by [1992Kim, 1992Shi,

1998Tak].

The (V1-xTix)Al3 (x = 0.01 to 0.25) alloys have been studied by [1988Uma]. The master alloys were

prepared in an argon plasma furnace, then these master alloys have been remelted in a high-purity alumina

crucible under an atmosphere of Ar-10 % H2 and annealed at 1000°C for 48 h. Optical and electron

microscopy, EMPA and XRD were used to study the structure of the alloys. The D022, TiAl3 type crystal

structure, was confirmed for the intermetallic phase with c/a changing from 2.202 to 2.210 in the studied

range.

Melt spinning was used to prepare Al-(Ti1-xVx)Al3 alloys studied by [1989Fra]. One Al-Ti and three

Al-Ti-V alloys containing 94 - 97 mass% Al were annealed at 200 and 600°C for 24 h and examined in as

received and annealed conditions. The fcc (Al) and bct (Ti1-xVx)Al3 with the D022 structure were identified

in all alloys. The V rich compounds were not observed.

The 8Ti-2.13V-Al (mass%) alloy prepared by mechanical alloying was studied by [1993Lee1]. The

obtained powder has been heated at 450°C for 10 h in a H2 atmosphere and were examined by XRD. The

crystal structure of (Ti0.8V0.2)Al3 was found to be bct, D022.

28


MSIT®

Al–Ti–V

The phase relations and the crystal structure of phases in the alloys based on (Ti1-xVx)Al3 were studied by

[1994Cha, 1995Cha, 1996Cha, 1998Cha]. The alloys were prepared by arc melting and chill cast in place

on a water-cooled Cu hearth. The compositions of the studied alloys were 25Ti-xV-(75-x)Al (x = 4, 6, 9),

10Ti-20V-70Al, 10Ti-20V-62Al, 10Ti-35V-55Al. The alloys were homogenized at 1250°C for 100 h

(25Ti-(75-x)Al-xV) or 1 h (the other alloys), cooled slowly to heat treatment temperatures in the range 800

- 1000°C and annealed for 3 to 100 h. Optical microscopy, SEM, TEM, XRD, electron diffraction were used

to examine the alloys. EDXS was used to carry out X-ray microanalysis of the precipitate phase and light

element analysis was performed using parallel electron energy loss spectrometer (PEELS) [1996Cha].

Phases with long period structures based on the L12 structure were found to be involved in phase

transformations during solidification and subsequent cooling of the alloys in the region between (Ti,V)Al3and (Ti,V)Al. In this region, the phases Ti2Al5 and Ti5Al11 were found and the existence of the

Ti2Al5+Ti5Al11+(Ti,V)Al3 equilibrium was suggested. The eutectic L V5Al8+(Ti,V)Al3 ( + ) was

observed in the chill cast 10Ti-70Al-20V and 10Ti-62Al-28V alloys, but after homogenization treatment at

1250°C for 1 h followed by furnace cooling to temperatures in the range of 800 to 1000°C and isothermal

aging, the + mixture has been transformed into + . The L + eutectic was suggested to be metastable,

the equilibrium eutectic was supposed to be L + . DTA was used by [1997Cha] to determine the liquidus

and solidus temperatures for the 10Ti-20V-70Al alloy.

The phase formation in Al-Ti-V alloys has been considered by [1993Cui, 1995Sha2].

Metastable phase diagram for the Ti-xV-3Al and Ti-xV-6Al (V up to 20 mass%) sections were determined

by [1985Kol2].

The ´´ martensitic transformation was studied in order to determine compositions of the alloys

sufficiently stable with respect to the ´´ martensite formation for transformation toughening of titanium

aluminide [1992Gru1, 1992Gru2, 1992Gru3, 1997Gru].

The Ti-4.0Al-15.4V and Ti-4.0Al-16.1V (mass%) shape memory (SM) alloys have been studied by means

of optical microscopy, TEM and XRD [1991Pak].

Site occupancies of V in TiAl ( ) and Ti3Al ( 2) alloys were analyzed using the atom location channeling

enhanced microanalysis (ALCHEMI) [1991Moh, 1999Hao]. A theoretical model for the sublattice site

occupancies and prediction of the stabilizing effects of V additions to the and 2 phases was presented by

[1999Yan]. The thermodynamic model was applied to predict the site substitution behavior in TiAl

[1990Nan]. The theoretical and experimental investigations of sublattice substitution of alloying elements,

including V, in TiAl and Ti3Al were summarized by [2000Yan].

The linear muffin tin orbital (LMTO) method was applied to calculate the electronic structure and total

energies of L10 ordered TiV2Al and Ti2VAl compositions from first principles [1993Ers].

The structural stability and cohesive properties of Ti2VAl as B2, D019 and orthorhombic phases have been

studied theoretically by [1999Rav]. The cluster variation method (CVM) was applied to calculate the / 2

phase equilibria [2001Kan]. The calculated isothermal sections of the Al-Ti-V phase diagram at 2100, 1800,

1500, 1200, 900 and 600 K were presented by [1989Kau]. Three stoichiometric compounds Ti3Al, TiAl and

TiAl3, and formation of the phase through the peritectoid reaction were accepted to exist in the binary

Al-Ti system used for ternary calculation.

Binary Systems

The Al-Ti phase diagram is accepted from [1993Oka, 2003Sch1], who used the assessed phased diagram

obtained by [1992Kat] from available experimental data and thermodynamic calculation. Recent data for

the TiAl-TiAl3 region by [2001Bra] shown by [2003Sch1] in addition to the version by [1992Kat] also are

taken into account.

The accepted Al-V system is given by [2003Sch2]. The Ti-V phase diagram is taken from [Mas2].

Solid Phases

Data on the solid phases pertinent to Al-Ti-V alloys are given in Table 1.

An existence of the wide range of the (Ti,V,Al) solid solutions (up to ~40-50 at.% Al at high temperatures)

and their stabilization with V additions were confirmed. The second order transformation of a disordered

29


MSIT®

Al–Ti–V

bcc (A2) phase into the ordered 0 phase with the CsCl (B2) structure was established [1994Ahm1,

1994Ahm2, 1995Sha2, 1996Sha]. The ordering temperature was found to decrease with increasing V

content at the constant Al content of 20 mass% and to increase with increasing Al content at the constant V

content of 20 mass% [1994Ahm1]. The lattice parameter of ( 0) phase was found to decrease with

increasing V and Al contents. [1994Ahm1] proposes the following equation to predict the lattice parameters

of the phase in the Al-Ti-V alloys:

a = 327 - 0.1x - 0.3y (pm),

where x and y are Al and V contents in at.% respectively.

The similar equation also was proposed by [1995Sha2] to define the lattice parameters of the and 0

phases:

a ( 0) = 327 - 0.154x - 0.236y (pm).

The Ti3Al based 2 phase extends to ~10 at.% V. Its lattice parameters decrease with increasing V and Al

contents. [1994Ahm1] proposes the following equations:

a 2= 589 - 0.4x - 0.2y

c 2 = 470 - 0.2x - 0.4y.

V dissolves in the phase up to ~20 at.% and results in decreasing lattice parameters of this phase

[1968Kor, 1986Has, 1988Has].

An existence of the continuous solid solutions (shown as in this assessment) between the isostructural

compounds TiAl3 and VAl3 (D022, TiAl3 type crystal structure) established earlier and considered by

[1993Hay, 1995Hay] was confirmed by subsequent investigations [1989Fra, 1991Par]. The continuous

solid solutions (Ti,V)Al3 were shown also by [1992Ahm] in the temperature range 1200 - 600°C. However,

later the same authors showed the continuous solid solutions only at temperatures below 1200°C, but at

temperatures 1200°C two different phases, (Ti,V)Al3 ( 1) on the TiAl3 base and (V,Ti)Al3 ( 2) based on

VAl3, were shown by [1994Ahm1, 1994Ahm2]. However, crystal structure data for these phases in the

ternary Al-Ti-V system were not reported and a difference between these phases is not clear.

An ordered superstructure of the (Ti,V)Al3 phase based on Ti8Al24 was observed at lower temperature, but

again the lattice parameters are given only for the binary phase by [1973Loo].

Data on the phases presented in ternary alloys of the region between and are contradictory similarly to

data for the alloys of the region between TiAl and TiAl3 in the binary Al-Ti system.

[1995Cha, 1996Cha, 1998Cha] observed besides TiAl3 a series of tetragonal one-dimensional (1d)

long-period superlattice, based on the L12 structure, or antiphase domain structures (APS), Ti5Al11 (D023)

and Ti2Al5 phases in the Ti25Al75-xVx (x = 4 to 9) alloys. The Ti25Al66V9 alloy was found to be almost

single-phase D023. The phases (Ti20V8Al72), possibly L12 (Ti28V5.4Al66.6, a = 394.7 pm) and TiAl2 were

reported by [1991Spa] in the 25Ti-67Al-8V alloy annealed at 1200°C for 16 h. The phases Ti5Al11+V5Al8were found in the alloy of the same composition obtained by sintering at 1150°C for 24 h by [1993Nak].

The phase composition + (TiAl2 based phase) was detected in the Ti-62.4Al-6.9V alloy in the

temperature range 900 - 600°C by [1992Ahm, 1994Ahm1].

The metastable phases ´, ´´, (disordered and ordered) were observed in ternary alloys quenched or aged

[1992Gru1, 1992Gru2, 1993Cui, 1995Sha2, 1996Sha]. The formation of phase was shown to obey the

electron concentration rule. The phase was shown to occur at 4.12 e/a. The characteristic value for Ti

alloys was obtained to be 4.223 [1993Cui]. The disordered phase was found to be formed from a

disordered bcc phase, while the ordered phase with a space group P3ml formed in the ordered 0 (B2)

phase with c ord = 2c disord [1995Sha2, 1996Sha]. Both the electron-to-atom ratio e/a and size factor affect

the stability of the phase [1995Sha2].

An ordered fct phase ( 1) arisen from the bcc phase was reported to transform into an ordered fco 1´

phase in the Ti-4.0Al-15.4V (mass%) alloy with a shape memory effect during quenching from 880°C in

ice water [1991Pak]. The lattice parameters of the phases were determined as follows: a = 323 pm;

a 1 = b 1 = 457 pm, c = 323 pm (b 1 = a , c 1 = a ); a 1´ = 457 pm, b 1´ = 490 pm, c 1´ = 299 pm.

The Al rich compounds of the Al-V binary system, V4Al23, V7Al45 and V2Al21, were not observed in

studied alloys of the Al rich region of the Al-Ti-V system.

[1996Sha, 1997Sha] reported a certain “new ternary phase” H (disordered) or H2 (ordered) with the crystal

structure (hP8, P63/mmc) fully consistent with this of 2 (Ti3Al) and the lattice parameters a = 558 1 pm,

2

30


MSIT®

Al–Ti–V

c = 450 1 pm. This phase was observed in the chill cast alloys Ti-28V-62Al and Ti-35V-55Al. The

composition of the H (H2) phase was not reported.


Data on invariant equilibria, which were suggested to take place in the ternary Al-Ti-V system by various

authors are summarized in Table 2. There are two main versions of the phase equilibria. The first version is

characterized by the equilibrium between and at high temperatures and + at lower ones. The second

version implies the + equilibrium up to melting.

There is no doubt about the existence of the invariant equilibrium L+ + though the temperature of this

equilibrium has not been determined experimentally but was just estimated by [1993Hay, 1995Hay] and

[1994Ahm1]. The + + phase field was observed at 1200°C by [1992Ahm, 1994Ahm1, 2000Kai] and at

1300°C by [2000Kai].

The invariant reaction L+ + was found to occur at ~1400°C from the L+ + and L+ + equilibria

[1970Vol] and accepted by [1993Hay, 1995Hay] at 1390°C. A coexistence of the and phases was

observed in the as cast alloys by [1994Ahm2].

The invariant reaction L+ + was proposed by [1993Hay, 1995Hay] from analysis of results obtained

in various earlier works. [1994Ahm1] has shown a similar reaction, L+ + 1, involving (Ti,V)Al3 ( 1)

phase with a crystal structure different from the D022, type TiAl3, ( ) phase.

The + + and + + ( 1) phase fields were observed down to 1200°C by [1992Ahm, 1994Ahm1] and

changed to + + and + + at lower temperatures. The + + and + + phase fields were observed also

by many other authors [1986Has, 1991Par, 1993Hay, 1995Hay]. So, the invariant reaction + + ( 1)

was suggested to take place [1993Hay, 1995Hay, 1992Ahm, 1994Ahm1]. However, [1965Kor, 1968Kor,

1970Vol, 1971Vol] have not detected any phase transformation in + alloys from ~1400 down to 550°C.

On a contrary, [1991Par] has not found the + equilibrium and has suggested the existence of the invariant

equilibria L+ + and L+ + followed by the + + and + + equilibria. The same equilibria are

shown in the calculated phase diagram by [1995Sha1]. The eutectic reaction L + has been considered to

be the equilibrium one by [1997Cha], that can be an evidence for the versions of [1991Par] and [1995Sha1].

A continuous solid solution between TiAl3 and VAl3 phases (D022) was found to exist from a study of alloys

annealed at temperatures below 1100°C [1956Jor, 1966Ram, 1989Fra, 1991Par, 1994Ahm1], but there is

no information on the structure of the as cast TiAl3-VAl3 alloys. Both aluminides are formed through the

peritectic reactions L+ and L+ in the binary systems Al-Ti and Al-V, respectively. The invariant

reaction L+ + in the Ti rich part of the TiAl3-VAl3 section was supposed by [1993Hay, 1995Hay,

1995Sha1]. The somewhat different reaction L+ + was proposed by [1991Par] from temperatures of

the reactions accepted in that work, l+ (1380°C), l+ (1350°C), in the Al-Ti binary system and the

invariant reaction in the ternary system estimated to be 1370°C. The invariant reaction L+ + 1 proposed

by [1994Ahm2] involves the 1 phase with a tetragonal crystal structure, based on TiAl3, which was not

given in detail. One more invariant reaction L+ 1+ 2 was proposed by [1994Ahm2] taking into account

an existence of two various phases for (Ti,V)Al3. However, the alloys along the TiAl3-VAl3 section have

not been studied by [1994Ahm1].

The + + (D023) phases were found by [1966Ram] in as cast Ti-5V-70Al alloy, but in the Ti-4V-67Al

alloy [1995Cha, 1996Cha, 1998Cha] has determined the (D022), (D023) and a phase with a long-period

superlattice close to Ti2Al5. It can be assumed that phase equilibria in this part of the ternary system are

more complicated than it was supposed earlier.

One of possible tentative versions of the reaction scheme in the ternary system is shown in Fig. 1.


There are four versions of the liquidus surface projection in literature [1991Par, 1993Hay, 1994Ahm2,

1995Hay, 1995Sha1]. The version by [1993Hay, 1995Hay] is assessed, that by [1995Sha1] is calculated.

DTA was used by [1991Par]; [1994Ahm2] studied only structures of as cast alloys. The fields of primary

crystallization of the , , , , and phases and four invariant liquid points related to invariant equilibria

listed in Table 2 are shown by [1991Par, 1993Hay, 1995Hay, 1995Sha1], but two separate fields (Ti,V)Al3

31


MSIT®

Al–Ti–V

phases, 1 and 2, and five invariant liquid points are shown by [1994Ahm2]. Figure 2 shows one more

version of the partial liquidus surface projection corresponding to the tentative reaction scheme in Fig. 1

and the accepted binary phase diagrams. The phase, summarizing various phases observed in a region

between and phases, and the single phase field are accepted. Data reported by [1991Par, 1992Cha1,

1992Cha2, 1994Ahm2, 1994Cha, 1997Cha] were used in addition to the data considered by [1993Hay,

1995Hay].

The phase fields of primary crystallization of Al rich phases V4Al23, V7Al25 and V2Al21 have not been

observed. Probably, their boundaries are very close to the Al-V side.

Isothermal Sections

The isothermal section at 1400°C was presented by [1993Hay, 1995Hay] from the critical consideration of

earlier results concerning the Ti rich part of the ternary phase diagram and the TiAl-V5Al8 section by

[1970Vol]. Taking into account DTA data by [1991Par] a wider liquid region and additionally the + +L

phase fields are shown in Fig. 3 of the present review.

Boundaries of the + phase field ( - tie lines) determined at 1300°C by [2000Kai] are shown in Fig. 4.

The isothermal section at 1200°C shown in Fig. 5 is based mainly on results by [1992Ahm, 1994Ahm1]. In

contrast to the similar section given by [1993Hay, 1995Hay] in accordance with [1970Vol], the TiAl-V5Al8section was not found to be pseudobinary at this temperature. A Ti solubility in V5Al8 determined by

[1992Ahm, 1994Ahm1] significantly exceeds that reported earlier [1993Hay, 1995Hay]. The existence of

ordered CsCl (B2) type 0 solid solutions was established [1994Ahm1]. Continuous solid solutions (TiAl3type) were shown by [1992Ahm] but two different (Ti,V)Al3 phases, 1 and 2, were given by [1994Ahm1].

Only (D022) phase is shown in Fig. 5 because there is no reason to suppose an existence of two phases

with different crystal structure in the TiAl3-VAl3 section at this temperature.

The isothermal section at 1100°C was presented by [1993Hay, 1995Hay] mainly from the results by

[1968Kor] taking into account earlier data on the Ti rich corner of the ternary phase diagram. The

TiAl-V5Al8 system was again supposed to be pseudobinary. However, the existence of the + + phase

equilibrium at 1100°C can not be quite reliably established taking into account that according to [1965Kor,

1968Kor, 1970Vol] this equilibrium was shown to exist down to 500°C but it was not confirmed at 1000°C

and lower temperatures by [1956Jor, 1986Has, 1991Par, 1992Ahm, 1994Ahm1, 1994Cha, 1997Cha]. So,

the phase equilbria at 1100°C are shown tentatively in Fig. 6.

The isothermal section at 1000°C is shown in Fig. 7. The section at this temperature was presented by

[1956Jor, 1986Has]. A Ti rich corner was given by [1969Tsu]. The earlier works were reviewed by

[1962Ere]. The assessed isothermal section at 1000°C was proposed by [1993Hay, 1995Hay]. The existence

of the continuous solid solutions between TiAl3 and VAl3 is in agreement with [1956Jor, 1986Has,

1988Uma]. A position of the 2+ + triangle has been determined also by [1992Cha1]. A more extended

2+ ( 0)+ field was shown by [1992Cha1, 1993Hay, 1995Hay]. The 2+ / boundary agrees well with

[1988Has]. Data on the structure of alloys located between the and regions, which were studied by

[1995Cha, 1996Cha, 1998Cha] were not reported for those annealed at 1000°C.

The isothermal section at 900°C was constructed from experimental [1991Par, 1992Ahm, 1994Ahm1] and

calculated data [1995Sha1, 1995Sha2]. There is a good agreement between all versions of the section.

Figure 8 shows the section based mainly on that by [1994Ahm1]. There is an agreement with the structure

of the Ti-20V-70Al and Ti-28V-62Al alloys ( + ) observed by [1994Cha, 1996Cha, 1998Cha].

The isothermal section at 800°C was determined by [1965Kor, 1968Kor, 1986Has, 1994Ahm1] and

assessed by [1993Hay, 1995Hay]. As it was noticed by [1993Hay, 1995Hay], the section proposed by

[1965Kor, 1968Kor] is unsuitable, because below the temperature of the invariant reaction + + , which

takes place above 900°C, the equilibrium + exists rather than + shown by [1965Kor, 1968Kor,

1970Vol] down to 500°C. The isothermal section at 800°C constructed by [1994Ahm1] was used to

represent that shown in Fig. 9.

The partial isothermal sections at 700 and 600°C given in accordance with the data by [1994Ahm1] are

shown in Figs. 10 and 11, respectively, taking into account the Ti-V phase diagram by [Mas2]. The binary

32


MSIT®

Al–Ti–V

Ti-V phase diagram [1981Mur] without the monotectoid reaction was used for the calculated section at

600°C by [2002Dip] (Fig. 12).

The isothermal sections at 2100, 1800, 1500, 1200, 900 and 600°C calculated by means of CALPHAD

method were presented by [1989Kau].


Various vertical sections given in earlier works ([1956Rau] (Ti rich part of sections at 2, 4, 6, 8 mass%),

[1969Tsu] (at 2, 4, 8 mass% V, 6, 8, 10 mass% Al), [1968Kor] (Ti-(V:Al = 3:1, 1:1, 1:3) up to 65, 85, 100

mass% (V+Al), respectively), [1970Vol] (TiAl-V5Al8), [1971And] (Ti3Al-V)) were constructed from

earlier versions of the Al-Ti phase diagram. So, they are not reproduced here.

Recently some calculated vertical sections were published [1995Ahm, 1995Sha1, 1997Sha]. Figure 13

gives the calculated section at 12 at.% V [1995Ahm]. The section at 50 at.% Ti [1997Sha] is shown in

Fig. 14. The similar section at 49 at.% Ti was given by [1995Sha1].


Aluminium and vanadium are the major alloying additions to titanium. The Ti-6Al-4V (mass%) alloy is the

most prevailing commonly used titanium alloy due to its superplasticity behavior and high specific strength.

Numerous publications concern this alloy, only a few of them will be pointed out in this review. Last 20

years alloys based on the TiAl and Ti3Al intermetallic compounds are considered as high temperature

materials for structural applications due to their low density, relatively high melting temperature, high

resistance to oxidation, good strength at elevated temperatures and good creep properties. However, they

show low ductility and fracture toughness at room temperature. A great number of investigations are

undertaken in order to improve the mechanical properties of the titanium aluminide by addition of other

elements, vanadium among them [1989Kim, 1992Nak, 1997Nak, 1999Flo]. Mechanical properties of V

containing TiAl based alloys are summarized by [1989Kim].

A positive temperature dependence of the yield stress with a peak temperature at 800°C was found for the

Ti-55Al-10V by [1990Wha]. A tensile and creep behavior of as cast Ti-48Al-3V alloy in different

microstructural conditions were investigated by [1992Naz]. This alloy was found to exhibit improved

strength and creep properties at room temperature compared with a nickel based superalloy. [1992Shi]

carried out compressive testing of + / 0 alloys in air-cooled and aged conditions and measured

microhardness of phases as well. It was shown that the 0 phase formed during air or water cooling from

950 to ~1200°C to room temperature (RT) gives rise to the high strength at RT. Upon aging the quenched

alloys at 550 ~750°C, a hard ordered 0 phase transformed to a ductile disordered phase with less Al

content. The specimens with higher amount of phase exhibit fracture strain higher than 35 % and the yield

stress not less than 700 MPa. The phase exhibits higher yield stress and a larger fracture strain than the 0

phase.

The effect of the control of microstructure on mechanical properties of a TiAl based alloy with 2 at.% V

was examined by [1993Has]. Mechanical properties, Young’s modulus, elongation, yield stress and fracture

stress at 77 K, 293 K and higher temperatures have been determined. It was concluded that the

microstructure is the most important parameter to improve ductility.

Fatigue tests of Ti-2V-48Al alloy were performed by cycling between Tmin = 100°C and Tmax ranging from

750 to 1400°C and a mechanism of thermal mechanical fatigue was proposed [1994Lee].

The addition of V to TiAl alloys was found to increase significantly yield stress from room temperature to

900°C. The maximum yield stress in Ti-10V-55Al alloy ( ) was observed to occur near 800°C (in single

crystal Ti-54Al in the range of 600 to 800°C) [1995Hah].

The dependencies of the tensile properties and creep resistance of (52Ti-48Al)-xV (x = 0 to 6) alloys were

studied by [2001Sun]. The increase of tensile strength and creep resistance with increasing V content was

attributed to the solid solution strengthening of this element in the phase. The appearance of 0 phase

deteriorated the creep resistance, room temperature strength and ductility.

The (Ti,V)Al3 compounds containing 1 to 25 at.% Ti as potential materials for the nuclear reactor

technology were deformed under compression at temperatures between 20 and 940°C [1988Uma]. A

33


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Al–Ti–V

remarkable improvement in ductility of VAl3 was found by the replacement of vanadium by titanium,

particularly at low temperature.

Resistivity, microhardness of the phase and hardness of alloys in the TiAl-V5Al8 section were studied by

[1970Vol] were shown to increase with V content in the alloys.

Magnetic susceptibility of the alloys in TiAl-VAl section has been determined by [1976Sid]. TiAl, VAl and

TixV1-xAl solid solutions were found to be weak paramagnetic in the temperature range from 78 to 660 K.

Thermophysical properties (heat capacity, thermal and temperature conductivity coefficients) of Ti alloys

containing up to 9 at.% Al and up to 3 at.% V (7 alloys) in the temperature range 750 to 1700 K were studied

by [1980Zar]. The following function was obtained for the density of solid solutions of the Al-Ti-V alloys:

d(x,y) = 4.544 - 0.0293x - 0.00852y+0.000988x2+0.0011751xy+0.0034009y2 (g cm-3),

where x, y are Al and V contents in at.%.

Superplasticity in Ti-2.5V-3Al and Ti-4V-3Al (mass%) alloys were studied by [2000Sal1] and [2000Sal2],

respectively.

High temperature ductility losses in + alloys (Ti-6Al-xV (x = 0 to 6 mass%)) were supposed to be related

to the lattice contraction during the transformation [1987Dam].

It was established that the most important factors controlling toughening of aluminide by dispersion of

the phase in Al-Ti-V alloys are thermodynamic stability of the phase with respect to martensitic

transformation, the martensite transformation volume change and the chemical compatibility of two

phases, which prevents chemical reaction at the / interface [1992Gru2]. It was found that the composition

region Ti-(45-55)V-(30-40)Al (mass%) is the range of the phase with an optimum combination of

transformation volume change and the chemical compatibility between the and phases. However, the

phase of this composition range was found not to have the required thermodynamic stability with respect to

the martensitic transformation during cooling to room temperature [1992Gru1]. A thermodynamic analysis

was performed to determine the optimum chemical composition of the phase possessing the required level

of stability with respect to martensitic transformation. It was found that this phase composition range is

within 10 mass% of the /( + ) phase boundary [1992Gru2].

The age-hardening response of a 0 Ti-39.5Al-21V alloy has been studied by [1995Ahm]. It was shown that

the hardness increase is due to precipitation, longer-time precipitation of 2 resulted in further hardness

increase. Increasing aging times resulted in a loss of coherency at the / 0 interfaces giving rise to a plateau

and a minimum in the aging curves at 600 and 700°C.

Miscellaneous

Vanadium atoms occupy Ti sites in Ti3Al [1999Hao, 2000Yan], preferentially substitute for Ti in TiAl

[1990Nan, 1991Moh]. According to [1999Hao, 2000Yan] V substitutes for both Ti and Al at low

concentration its probability to occupy the Ti sites increases with increasing V content. V atoms substitute

for both sites in TiAl [1988Has, 2001Kan]. V atoms replace for Ti atoms in TiAl3 crystal lattice [1989Fra,

1990Abd].

The TiAl3 based aluminides obtained from dilute Al-(Ti,V) melts have shown the crystal structure similar

to that of TiAl3 and compositions, which agrees with the formula (Ti1-xVx)Al3, and disagrees with the

existence of a homogeneity range of the phase on the Al rich side up to ~14 at.% V [1990Abd]. Lattice

parameters of (Ti1-xVx)Al3 were found to decrease linearly with increasing V content [1989Fra], the c/a

ratio decreased from 2.210 to 2.201 [1988Uma].

The calculation of the electronic structure and the total energy of Ti2VAl using the self-consistent

tight-binding linear muffin-tin orbital method has predicted that Ti2VAl is more stable in the B2 phase than

in D019 [1999Rav].

The metastable diagrams observed from quenched alloys Ti-3Al-V and Ti-6Al-V (mass%) are shown in

Figs. 15 and 16, respectively. Figure 17 shows the nature of different phases obtained from quenched

Al-Ti-V alloys in the Ti-rich corner [1985Kol2]. [1993Cui] studied the phase formation, which was

shown to obey the electron concentration rule. The phase was observed to occur at electron concentration

value of ~4.12 in alloys from Ti-12V (at.%) to Ti-14V-3Al.

34


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Al–Ti–V

[1995Sha2] observed formation of a disordered diffuse phase in a solidified disordered phase and an

ordered phase (space group P3m1) in the ordered B2 phase. The addition of V destabilizes the high

temperature phase and enhances the formation. The ordered phase was observed in a melt-spun

Al-Ti-V ribbons [1996Sha]. The phase transformation sequence in the Ti-10V-40Al alloy was determined

to be L L+ + +B2 2+ +B2 2+ +B2+ ord [1995Sha1, 1997Sha].

[1991Pak] investigated shape memory Al-Ti-V alloys. It was found that a martensitic transformation took

place in the Ti-15.4V-4.0Al alloy during quenching. The transformation resulted in volume change of

0.2 %. This small volume change and atomic ordering were considered to be an origin for the shape memory

effect in Ti based alloys. There was no evidence of the presence of phase in the quenched Ti-15.4V-4Al

alloy, but it was observed in the Ti-16.1V-4Al. It was supposed that atomic ordering suppresses the phase

formation.

The formation of the L12 type compound in the Ti-(75-x)Al-xV (x up to 8) alloys, arc melted and annealed

at 1200°C for 16 h, was observed by [1991Spa]. The phases L12 (Ti-66.6Al-5.4V), D022 (Ti-72Al-8V) and

an insignificant amount of the phase with composition of Ti-12V-65Al were determined in the Ti-67Al-8V

alloy. The amount of L12 phase increased with V concentration. The phase composition D023+V5Al8 was

found in the alloys of the same composition sintered at 1150°C for 24 h by [1993Nak].

A small amount of precipitates of a Al-O-Ti ternary compound was observed in single-phase alloy Ti-5V-54Al.

Such compound was not found in the two-phase Ti-45Al-5V alloy because of the oxygen scavenging effect of the

2 phase [2001Cao, 2002Cao]. The crystal structure of this compound was reported to be cubic (a = 690 pm)

[2001Cao], but later it was refined and found to be triclinic with a = 490, b = 770, c = 940 pm, = 105.37°, =

75.01°, = 93.16° [2002Cao].

Hydrogenation of Ti75-xVxAl25 (x = 0, 15, 25) was studied by [2001Ish]. The addition of V resulted in

reduced hydrogen capacity in comparison with Ti3Al. Nevertheless, the alloys containing V absorbed

0.4-0.1 H/M (1 - 2 mass%). The 50 % desorption temperature increased by alloying with V. The bcc phase

was obtained in the alloy containing 25 at.% V. The alloy with 15 at.% V turned into amorphous.

Compositions of the and phases in the Ti-4V-6Al alloy annealed in industrial conditions were found to

be different from equilibrium. So, commercial alloys are metastable [1985Kol1].

The texture and microstructure in Ti-4V-6Al formed during hot rolling at temperatures between 750 and

1050°C were investigated by [1990Ina]. [1993Lee2] has considered variation of equilibrium fraction of the

phase in Ti-4V-6Al versus temperature (calculation and experimental data). Grain growth kinetics at

temperatures in the range 1050 - 1200°C was studied by [2000Gil].

Isothermal hot compressing tests were carried out in order to estimate the hot deformation mechanisms in

extra-low interstitial (ELI) grade Ti-4V-6Al for optimization the workability of the alloys in + phase

fields [2000Ses]. The transus of this alloy was about 975°C.

The kinetics of + transformation in Ti-4V-6Al was studied by means of electrical resistivity technique

[2001Mal1]. The phase composition of the alloy has been studied in the temperature range 950 to 750°C

and the phase equilibria were calculated using ThemoCalc. A good agreement was found between

experimental and calculated phase compositions. Differential scanning calorimetry (DSC) was used for

investigation of transformation in Ti-4V-6Al by [2001Mal2]. Continuous cooling transformation

(CCT) diagrams were calculated.

Phase transformation kinetics in the Ti-2V-6Al and Ti-6V-6Al (mass%) alloys were studied under

continuous cooling conditions [1988Dam]. The transformation in the alloy with 2V occurred more rapidly

than that in the alloy with 6V. For the cooling rate of 10°C s-1 the transformation onset temperatures

were found to be 865 and 735°C for the alloys with 2V and 6V, respectively.

It should be noted that there are contradictory data on the phase composition of the alloys in the

TiAl2-TiAl3-VAl3-V5Al8 region at the temperatures higher than ~1100°C. Some of these data allow to

suppose an existence of the phase equilibria + + , + + after crystallization and a number of phase

transformations at lower temperatures. However, the data available are insufficient to confirm this version.

35


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Al–Ti–V

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Al–Ti–V

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Al–Ti–V

[1992Cha1] Chaudhury, P.K., Rack, H.J., “Ti-Al-V Ternary Phase Stability at Elevated Temperatures”,

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[1992Shi] Shi, J.-D., Pu, Z., Zhong, Z., Zou, D., “Improving the Ductility of (TiAl) Based Alloy by

Introducing Disordered Phase”, Scr. Metall. Mater., 27, 1331-1336 (1992) (Crys.

Structure, Equi. Diagram, Experimental, 5)

[1993Cui] Cui, Y., Li, D., Wan, X., “ Phase Formation in Ti Alloys” (in Chinese), Acta Metall.

Sinica, 29, A61-A67 (1993) (Equi. Diagram, Experimental, 9)

[1993Ers] Erschbaumer, H., Podloucky, R., Rogl, P., Temnitschka, G., Wagner, R., “Atomic

Modelling of Nb, V, Cr and Mn Substitutions in -TiAl. I: c/a Ratio and Site Preference”,

Intermetallics, 1, 99-106 (1993) (Crys. Structure, Theory, 31)

[1993Has] Hashimoto, K., Kimura, M., “Effects of Third Element Additions on Mechanical Properties

of TiAl”, “Structural Intermetallics”, Darolia, R., Lewandowski, J.J., Liu, C.T., Martin,

P.L., Miracle, D.B., Nathal, M.V. (Eds.), Miner. Met. Mater. Soc., 309-318 (1993) (Equi.


[1993Hay] Hayes, F.H., “Aluminium - Titanium - Vanadium”, MSIT Ternary Evaluation Program, in


GmbH, Stuttgart; Document ID: 10.13932.1.20, (1993) (Crys. Structure, Equi. Diagram,

Assessment, 42)

[1993Lee1] Lee, K.-M., Lee, J.-H., Moon, I.-H., “Effects of V and Zr Addition on Lattice Parameters of

Al3Ti Phase in Mechanically Alloyed Al-8 wt.% Ti Alloys”, Scr. Metall. Mater., 29,


[1993Lee2] Lee, H.M., Soh, J.-R., Lee, Z.-H., Kim, Y.-S., “Effect of Alloy Composition on the Volume

Fraction of Beta Phase in Duplex Titanium Alloys”, Scr. Metall. Mater., 29, 497-501 (1993)

(Experimental, 35)

[1993Nak] Nakayama, Y., Mabuchi, H., “Formation of Ternary L12 Compounds in Al3Ti-Base

Alloys”, Intermetallics, 1, 41-48 (1993) (Crys. Structure, Experimental, 40)

[1993Oka] Okamoto, H., “Al-Ti (Aluminium - Titanium)”, J. Phase Equilibria, 14, 120-121 (1993)

(Crys. Structure, Equi. Diagram, Review, #, 16)

[1994Ahm1] Ahmed, T., Flower, H.M., “Partial Isothermal Sections of Ti-Al-V Ternary Diagram”,

Mater. Sci. Technol., 10, 272-288 (1994) (Equi. Diagram, Experimental, Review, #, *, 40)

38


MSIT®

Al–Ti–V

[1994Ahm2] Ahmed, T., Rack, H.J., Flower, H.M., “Liquidus Projection of Ti-Al-V System Based on

Arc Melted and Cast Microstructures”, Mater. Sci. Technol., 10, 681-690 (1994) (Equi.

Diagram, Experimental, Review, #, *, 26)

[1994Bra] Braun, J., Ellner, M., Predel, B., “On the Structure of High-Temperature Ti1-xAl1+x Phase”

(in German), J. Alloys Compd., 203, 189-193 (1994) (Crys. Structure, Experimental, 8)

[1994Cha] Chang, W.-S., Muddle, B.C., “Phase Relations in an Al3(Ti,V) Based in-situ Composite”,

Micron, 25, 519-525 (1994) (Crys. Structure, Equi. Diagram, Experimental, 23)

[1994Lee] Lee, E.U., “Thermal-Mechanical Fatigue of Ti-48Al-2V Alloy and its Composite”, Metall.

Mater. Trans., 25A, 2207-2212 (1994) (Experimental, 24)

[1995Ahm] Ahmed, T., Hayes, F.H., Rack, H.J., “Age-Hardening Response of 2 Ti-Al-V”, Mater. Sci.

Eng., A192, 155-164 (1995) (Crys. Structure, Equi. Diagram, Theory, #, 13)

[1995Cha] Chang, W., Muddle, B.C., “Microstructure in Chill-Cast Al3Ti-Based Al-Ti-V Ternary

Alloys”, Mater. Sci. Eng., A192/193, 233-239 (1995) (Crys. Structure, Equi. Diagram,

Experimental, 18)

[1995Hah] Hahn, Y.D., Whang, S.H., “Deformation and Microstructure in L10 Type Ti-Al-V Alloys”,

Metall. Mater. Trans., 26A, 113-131 (1995) (Crys. Structure, Experimental, 68)

[1995Hay] Hayes, F.H., “The Al-Ti-V (Aluminum - Titanium - Vanadium) System”, J. Phase

Equilibria, 16, 163-176 (1995) (Crys. Structure, Equi. Diagram, Review, 42)

[1995Sha1] Shao, G., Tsakiropoulos, P., Miodovnik, A.P., “Phase Transformations in Ti-40Al-10V”,

Intermetallics, 3, 315-325 (1995) (Crys. Structure, Equi. Diagram, Experimental,

Theory, #, 16)

[1995Sha2] Shao, G., Miodovnik, A.P., Tsakiropoulos, P., “ -Phase Formation in V-Al and Ti-Al-V

Alloys”, Phil. Mag., 71, 1389-1408 (1995) (Crys. Structure, Equi. Diagram,

Experimental, 35)

[1995Sha3] Shao, G., Tsakiropoulos, P., Miodovnik, A.P., “The Lamellar + Structure in

Al-30Ti-20V Alloy”, Scr. Metall. Mater., 33, 13-17 (1995) (Crys. Structure, Equi. Diagram,

Experimental, 11)

[1996Cha] Chang, W-S., Muddle, B.C., “Precipitation of (Ti,V)2Al(C,N) in Multiphase Al-Ti-V

Alloys”, Mater. Sci. Eng., A207, 64-71 (1996) (Crys. Structure, Equi. Diagram,

Experimental, 18)

[1996Sha] Shao, G., Tsarikopoulos, P., Miodownik, A.P. “Ordering and Decomposition of the Phase

in Melt-Spun TiAl1-xVx Alloys”, Mater. Sci. Eng., A216, 1-10 (1996) (Equi. Diagram,

Experimental, 15)

[1996Tre] Tretyachenko, L.A., “On the Ti-Al System”, Fifth International School “Phase Diagrams

in Material Science”, Katsyveli, Crimea, Ukraine, Abstracts, Sept. 23-29, 118 (1996) (Equi.


[1997Bul] Bulanova, M., Tretyachenko, L., Golovkova, M., “Phase Equilibria in Ti-Rich Corner of the

Ti-Si-Al System”, Z. Metallkd., 88, 256-267 (1997) (Crys. Structure, Equi. Diagram,

Experimental, 15)

[1997Cha] Chang, W.-S., Muddle, B.C., “Microstructure and Properties of Duplex

-Al3(Ti,V)/ -(Ti,V) Alloys”, Metall. Mater. Trans., 28A, 297-238 (1997) (Crys. Structure,

Equi. Diagram, Experimental, 38)

[1997Gru] Grujicic, M., Dang, P., “Martensitic Transformation in a Dispersed Ti-Al-V-Fe -Phase and

its Effect on Fracture Toughness of -Titanium Aluminide”, Mater. Sci. Eng., A224,


[1997Nak] Naka, S., Khan, T., “Designing Novel Multicomponent Intermetallics: Contribution of

Modern Alloy Theory in Developing Engineering Materials”, J. Phase Equilib., 18,

635-649 (1997) (Review, 17)

[1997Sha] Shao, G., Tsakiropoulos, P., “Ultra-Thin Twin Plates and Growth Domains in the Phase

as a Product of the B2 Phase Decomposition in Ti-40 at.% Al-10 at.% V”, Philos. Mag. A,

75A, 657-676 (1997) (Crys. Structure, Equi. Diagram, Experimental, 19)

39


MSIT®

Al–Ti–V

[1998Cha] Chang, W.-S., Muddle, B.C., ”Intermediate Phases and Phase Relations in the Composition

Range Al3(Ti,V) to TiAl in the Al-Ti-V system”, Mater. Sci. Eng., A251, 232-242 (1998)


[1998Tak] Takeyama, M., Ohmura, Y., Kikuchi, M., Matsuo, T., “Phase Equilibria and Microstructural

Control of Gamma TiAl Based Alloys”, Intermetallics, 6, 643-646, (1998) (Equi. Diagram,

Theory, 20)

[1999Flo] Flower, H.M., Christodoulou, J., “Phase Equilibria and Transformations in Titanium

Aluminides”, Mater. Sci. Technol., 15, 45-52 (Crys. Structure, Equi. Diagram, Review, 46)


Elements in TiAl and Ti3Al Alloys”, Acta Mater., 47, 1129-1139 (1999) (Crys. Structure,

Experimental, 41)

[1999Rav] Ravi, C., Vajeeston, P., Mathijaya, S., Asokamani, R., “Electronic Structure, Phase

Stability, and Cohesive Properties of Ti2XAl (X = Nb, V, Zr)”, Phys. Rev. B. 60,

15683-15690 (1999) (Crys. Structure, Theory, 32)

[1999Yan] Yang, R., Hao, Y.L., “Estimation of ( + 2) Equilibrium in Two Phase Ti-Al-X Alloys by

Means of Sublattice Site Occupancies of X in TiAl and Ti3Al”, Scr. Mater., 41, 341-346

(1999) (Equi. Diagram, Theory, 13)

[2000Gil] Gil, F.J., Planell, J.A., “Grain Growth Kinetic of the Near Alpha Titanium Alloys”, J.

Mater. Sci. Lett., 19, 2023-2024 (2000) (Experimental, 13)

[2000Kai] Kainuma, R., Fujita, Y., Mitsui, H., Ohnuma, I., Ishida, K., “Phase Equilibria around

(hcp), (bcc) and (L10) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855-867

(2000) (Equi. Diagram, Experimental, #, 29)

[2000Sal1] Salam, A., Hammond, C., “Superplasticity in Ti-3Al-2.5V”, J. Mater. Sci. Lett., 19,

1731-1733 (2000) (Experimental, 12)

[2000Sal2] Salam, A., Hammond, C., “Activation Energy for Superplastic Flow in Ti-3Al-4V Alloy”,

J. Mater. Sci. Lett., 19, 2155-2156 (2000) (Experimental, 6)

[2000Ses] Seshacharyulu, T., Medeiros, S.C., Frazier, W.G., Prasad, Y.V.R.K., “Mechanisms of Hot

Working in Extra-Low Interstitial Grade Ti-6Al-4V with Equiaxed ( + ) Microstructure”,

Z. Metallkd., 91, 475-480 (2000) (Experimental, 14)

[2000Yan] Yang, R., Hao, Y., Song, Y., Guo, Z.-X., “Site Occupancy of Alloying Additions in

Titanium Aluminides and Its Application to Phase Equilibrium Evaluation”, Z. Metallkd.,

91, 296-301 (2000) (Crys. Structure, Theory, Review, 38)


Binary System Ti-Al”, Metall. Mater. Trans., 32A, 1037-1047 (2001) (Crys. Structure,

Equi. Diagram, Experimental, #, 34)

[2001Cao] Cao, G.H., Liu, Z.G., Shen, G.J., Liu, J.-M., “Identification of a Cubic Precipitate in

-Titanium Aluminides”, J. Alloys Compd., 325, 263-268 (2001) (Crys. Structure,

Experimental, 16)

[2001Ish] Ishikawa, K., Hashi, K., Suzuki, K., Aoki, K., “Effect of Substitutional Elements on the

Hydrogen Absorption-Desorption Properties of Ti3Al Compounds”, J. Alloys Compd., 314,


[2001Kan] Kang, S.-Y., Onodera, H., “Analyses of HCP/D019 and D019/L10 Phase Boundaries in

Ti-Al-X (X = V, Mn, Nb, Cr, Mo, Ni and Co) Systems by the Cluster Variation Method”,

J. Phase Equilib., 22, 424-430 (2001) (Equi. Diagram, Theory, 15)

[2001Mal1] Malinov, S., Guo, Z., Sha, W., Wilson, A., “Differential Scanning Calorimetry Study and

Computer Modeling of Phase Transformation a Ti-6Al-4V Alloy”, Metall. Mater.

Trans., 32A, 879-887 (2001) (Equi. Diagram, Experimental, Theory, 34)

[2001Mal2] Malinov, S., Markovsky, P., Sha, W., Guo, Z., “Resistivity Study and Computer Modeling

of the Isothermal Transformation Kinetics of Ti-6Al-4V and Ti-6Al-2Sn-4Zr-2Mo-0.08Si

Alloys”, J. Alloys Compd., 314, 181-192 (2001) (Equi. Diagram, Experimental, Theory, 16)

40


MSIT®

Al–Ti–V

[2001Sun] Sun, F.-S., Cao, C.-X., Kim, S.-E., Lee, Y.-T., Yan, M.-G., “Alloying Mechanism of Beta

Stabilizers in a TiAl Alloy”, Metall. Mater. Trans., 32A, 1573-1589 (2001) (Crys. Structure,


[2002Cao] Cao, G.H., Liu, Z.G., Shen, G.J., Liu, J.-M., “Oxide Precipitation in V-Doped TiAl-Based

Alloys”, Mater. Sci. Eng., 328, 177-180 (2002) (Crys. Structure, Equi. Diagram,

Experimental, 17)

[2002Dip] Diplas, S., Tsakiropoulos, P., Shao, G., Watts, J.F., Matthew, J.A.D., “A Study of Alloying

Behavior in the Ti-Al-V System”, Acta Mater., 50, 1951-1960 (2002) (Equi. Diagram,

Experimental, Theory, #, 34)

[2003Kar] Karpets, M.V., Milman, Yu.V., Barabash, O.M., Korzhova, N.P., Senkov, O.N., Miracle,

D.B., Legkaya, T.N., Voskoboynik, I.V., “The Influence of Zr Alloying on the Structure and

Properties of Al3Ti”, Intermetallics, 11, 241-249 (2003) (Crys. Structure, Equi. Diagram,

Experimental, 16)

[2003Sch1] Schmid-Fetzer, R., “Al-Ti (Aluminium - Titanium)”, MSIT Binary Evaluation Program,

MSI, Stuttgart (2003) (Crys. Structure, Equi. Diagram, Thermodyn., Review, #, 85)

[2003Sch2] Schuster, J.C., “Al-V (Aluminium - Vanadium)”, MSIT Binary Evaluation Program, MSI,

Stuttgart (2003) (Crys. Structure, Equi. Diagram, Review, #, 36)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 664.2

Al

< 660.452

cF4

Fm3m

Cu

a = 404.96

0 to 0.6 at.% Ti [2003Sch1]

0 to 0.2 at.% V [2003Sch2]

[V-C2]

pure Al at 25°C [Mas2]

, (Ti1-xVx)

( Ti)(h)

1670 - 882

(V)

< 1910

cI2

Im3m

W

a = 330.65

a = 302.38

(A2) [V-C2]

0 x 1 at 882°C [Mas2, 1981Mur]

congruent melting at 1605°C at x = 0.32

[Mas2]

dissolves up to 44.8 at.% Al at 1490°C

[1993Oka, 2003Sch1], ~46 at.% Al at

1520°C [1996Tre, 1997Bul] at x = 0

dissolves up to ~50 at.% Al at ~1670°C,

x = 1 [Mas2]

pure Ti [Mas2]

pure V at 25°C [1981Kin, Mas2]

* 0, Ti1-x-yVxAly cP2

Pm3m

CsCl

a = 318

B2 ordered form of the solid solutions

for the 0 phases (Ti-26V-43Al,

Ti-26V-42Al, Ti-30V-39Al) air cooled

from 1150°C [1992Shi]

41


MSIT®

Al–Ti–V

, (Ti1-x-yVxAly)

( Ti)(r)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

[Mas2, V-C2]

y = 0.473 at 1490°C [1993Oka,

2003Sch1], y 0.48 at 1520°C

[1996Tre, 1997Bul] at x = 0

x =~0.03 at y = 0 [Mas2]

pure Ti at 25°C [1981Kin]

2, (Ti1-xVx)3-yAly

Ti3Al

< 1164

hP8

P63/mmc

Ni3Sn

a = 579.2

c = 462.9

a = 580.6

c = 465.5

a = 574.4

c = 462.4

a = 576 to 572

c = 464 to 460

a = 579

c = 465 to 398

D019 ordered phase [V-C2]

~20 to 39.6 at.% Al, maximum at 30.9

at.% Al [1993Oka, 2003Sch1]

[V-C]

at 22 at.% Al [L-B]

at 38 at.% Al [L-B]

in binary Al-Ti alloys [1994Ahm1]

in the Al-Ti-V system [1986Has]

, TiAl

< 1463

tP4

P4/mmm

CuAu

a = 400.5

c = 407.0

a = 400.0 0.1

c = 407.5 0.1

a = 398.4 0.1

c = 406.0 0.1

L10 ordered phase, 46.7 to 66.5 at.% Al

[1993Oka, 2003Sch1], 50 to 62 at.% Al

at 1200°C [2001Bra]

[V-C]

at 50 at.% Al [2001Bra]

at 62 at.% Al [2001Bra]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

42


MSIT®

Al–Ti–V

, TiAl2< 1199 tP4

P4/mmm

CuAu

oC12

Cmcm

ZrGa2

tI24

I41/amd

HfGa2

a = 403.0

c = 395.5

a = 1208.84

b = 394.61

c = 402.95

a = 397.0

c = 2430.9

summarizes several phases [2003Sch1]:

Ti1-xAl1+x, 63 to 65 at.% Al at 1300°C,

stable range 1445 - 1170°C [2001Bra],

listed as orthorhombic, Pmmm with

pseudotetragonal cell,

range ~1445 - 1424°C [1990Sch]

for Ti36Al64 at 1300°C [2001Bra]

metastable modification of TiAl2,

observed only in as cast alloys

[2001Bra], listed as TiAl2(h) (66 to 67

at.% Al, 1433 - 1214°C) [1990Sch]

stable structure of TiAl2 < 1216°C

[2001Bra]

66 to 67 at.% Al at 1000°C [2001Bra]

listed as TiAl2(r) by [1990Sch]

[2001Bra]

stable structure of TiAl2 of T < 1216°C

[2001Bra]

tetragonal

superstructure of

CuAu type

tI16

I4/mmm

ZrAl3tP28

I4/mmm

Ti2Al5

* a = 395.5

* c = 403.0

a 389

c 1620

a = 390.53

c = 2919.63


Ti5Al11, stable range 1416 - 995°C

[2001Bra], 66 to 71 at.% Al at 1300°C

[2001Bra] (including the stoichiometry

Ti2Al5)

at 64 at.% Al [1994Bra]

in the Ti25Al75-xVx (x = 4 to 9) alloys

[1995Cha, 1998Cha];

in the 10Ti-55Al-35V alloy annealed at

1250°C/1 h+900°C/24 h, WQ

[1996Cha]

, (Ti1-xVx)Al3

TiAl3(h)

< 1393

VAl3 1270

tI8

I4/mmm

TiAl3(h)

a = 384.88

c = 859.82

a = 378.1 0.1

c = 831.5 0.6

D022 ordered phase [V-C]

0 x 1 [1956Jor, 1966Ram, 1986Has,

1989Fra]

74.2 to 75 at.% Al [1993Oka, 2003Sch1]

1387 - 735°C; [2003Kar]

[2003Sch2]

TiAl3(l)

< 950 (Ti rich)

< 735 (Al rich)

tI32

I4/mmm

TiAl3(l)

a = 387.7

c = 3382.8

74.5 to 75 at.% V [2001Bra]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

43


MSIT®

Al–Ti–V

, V5Al8 1408

cI52

I43m

Cu5Zn8

a = 923.45

a = 925.3

[1989Mur, 2003Sch2]

in the chill cast Ti-20V-70Al alloy

[1994Cha];

dissolves up to ~12.% Ti at 1200°C, < 1

at.% Ti at temperatures below 900°C

[1992Ahm, 1994Ahm1], 1 at.% V at

900°C [1991Par]; up to 5 mass% Ti at

1000°C, 1 mass% Ti at 800°C

[1986Has]

V4Al23

< 736

hP54

P63/mmc

V4Al23

a = 769.28

c = 1704.0

[1989Mur, 2003Sch2]

V7Al45

< 730

mC104

C2/m

V7Al45

a = 2563

b = 763.7

c = 1108.8

=128.83°

[1989Mur, 2003Sch2]

V2Al21

< 690

cF184

Fd3m

VAl10

a = 1449.2 [1989Mur, 2003Sch2]

´ hP2

P63/mmc

Mg

metastable phase, 0 to 5 at.% V

in the Ti-V system [1987Mur]

in Ti-(0-8)V-3Al,

Ti-6Al-(0-7.5)V (mass%) [1985Kol2]

´´ oC4

Cmcm

U

a = 490

b = 457

c = 299

metastable phase, 5 to 15 at.% V

in the Ti-V system [1987Mur]

in Ti -V-(3, 6)Al (mass%) alloys

[1985Kol2]

in Ti -(45-55)V-(30-40)Al (mass%)

[1992Gru1, 1992Gru2, 1992Gru3,

1993Cui, 1997Gru]

for as quenched Ti-15.4V-4.0Al alloy

with a a = 323 [1991Pak]

T hP3

P6/mmm

CrTi

or P3m1

a = 460

c = 282

A20, metastable phase, 11 to ~50 at.% V

in the Ti-V system [1987Mur];

in Al-Ti-V alloys [1993Cui, 1995Sha2,

1996Sha]

in Ti-V-(3, 6)Al (mass%) alloys

[1985Kol2]

[1991Pak]

TiAl3 cP4

P3m1

Cu3Au

a = 397.2 metastable phase obtained from splat

cooling at 85 at.% Al [2001Bra]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

44


MSIT®

Al–Ti–V

Table 2: Invariant Reactions Detected in the Ternary Al-Ti-V System

[1993Hay, 1995Hay] [1991Par] [1994Ahm2] [1995Sha1]

L + +

1460°C > T > 1400°C

L + + L + +

1450°C > T > 1400°C

L + +

L + +

1390°C

(~1400°C [1970Vol])

L + +

1390°C

L + +

1387°C > T > 1360°C

L+ +

~1370°C

L + + 1

1395°C > T > 1360°C

L + “ ” +

L + + L + + 1

1360°C > T > 1300°C

L + + L + +

L + +

1320°C > T > 1215°C

L + +

L + 1 + 2

1360°C > T > 1300°C

+ +

1175°C > T > 990°C

2+ +

1118°C > T > 1100°C2 + 0 +

1200°C > T > 900°C

+ +

1100°C > T > 1000°C

+ +

1200°C > T > 900°C

[1994Ahm1]

+ + 1

[1994Ahm1]

45


MSIT®

Al–Ti–V

Fig

. 1:

A

l-T

i-V

. R

eact

ion s

chem

e

Al-

Ti

Ti-

VA

-B-C

βα

+ β

´

67

5e 4

Al-

Ti-

VA

l-V

l +

βδ

14

08

p4

β +

l

α1

490

p1

α +

l

γ1

463

p2

L +

α

β +

γ1463>

T>

1390

U1

L +

βγ

+δ

13

90

U2

γ +

l

ξ1

416

p3

ξ +

l

ε1

393

p5

L +

δγ+

ε1

280

U4

l +

δε

12

70

p6

αα 2

+β/

β 0

e 1

γ +

ξη

11

99

p7

αα 2

+γ

11

18

e 2

ξη

+ ε

99

0e 3

L +

ξγ

+ε

1393>

T>

1260

U3

γ +

δβ

+ ε

T>

1000

U5

αα 2

+β 0

+γ

11

00

E

γ +

ξε

+η

U6

α +

β0

α 2 +

β´675>

T

L+

β+γ

L+

γ+δ

L+

γ+ε

γ+δ+

ε

β+γ+

δ

ε+η+

ξγ+

ε+η

β+γ+

εβ+

δ+ε

α+β+

γ γ+ε+

ξ

α 2+

β 0+

γα+

α 2+

β/β 0

α+α 2

+β´

α 2+

β´+

β 0

46


MSIT®

Al–Ti–V

20

40

60

80

20 40 60 80

20

40

60

80

Ti V

Al Data / Grid: at.%

Axes: at.%

p4

αβ

U1

U2

U3 U

4

γ

δ

ε

ξ

p6

p5

p3

p2

p1

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%

L

α

γ

β

δ

ξ

Fig. 2: Al-Ti-V.

Tentative partial

liquidus projection

Fig. 3: Al-Ti-V.

Tentative isotermal

section at 1400°C

47


MSIT®

Al–Ti–V

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%

β0

β

α

γ

δ

ε

L

ξVAl

3TiAl

3

30

40

50

60

70

10 20 30 40 50

30

40

50

60

70

Ti 80.00

V 0.00

Al 20.00

Ti 20.00

V 60.00

Al 20.00

Ti 20.00

V 0.00

Al 80.00Data / Grid: at.%

Axes: at.%

γ

α

β

Fig. 5: Al-Ti-V.

Partial isotermal

section at 1200°C

Fig. 4: Al-Ti-V.

The region of the ,and phases at

1300°C

48


MSIT®

Al–Ti–V

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%

TiAl3 VAl

3

L

β

β0

α

α2

γ

η

ξ ε

δ

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%

L

ββ0

α

α2

γ

δ

ε VAl3

TiAl3

ξ

η

Fig. 6: Al-Ti-V.

Tentative partial

isotermal section at

1100°C

Fig. 7: Al-Ti-V.

Partial isothermal

section at 1000°C

49


MSIT®

Al–Ti–V

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%

L

α

α2

β0

β

γ

η

TiAl3 VAl

3

δ

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%

L

εTiAl

3 VAl3

δη

γ

α2

α

β0

β

Fig. 8: Al-Ti-V.

Partial isothermal

section at 900°C

Fig. 9: Al-Ti-V.

Partial isothermal

section at 800°C

50


MSIT®

Al–Ti–V

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%

L

TiAl3 VAl

3

η

ε

γ

α2

α

β0

β

δ

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%(Al)

ε

η

δ

TiAl3 VAl

3

γ

α2

αα+β

1

β0

ββ

1+β

2 β2

Fig. 10: Al-Ti-V.

Partial isothermal

section at 700°C

Fig. 11: Al-Ti-V.

Partial isothermal

section at 600°C

51


MSIT®

Al–Ti–V

20

40

60

80

20 40 60 80

20

40

60

80

Ti V


Axes: at.%

εTiAl3 VAl

3

β0

βα

α2

γ

δη

10 20 30 40

750

1000

1250

1500

1750

Al 0.00

Ti 79.00

V 21.00

Al 50.00

Ti 29.00

V 21.00Al, at.%

Te

mp

era

ture

, °C

L

β

α+β

α2+β/β0

β0

α+β0

γ+β/β0

α2+γ+β0

Fig. 12: Al-Ti-V.

Calculated isothermal

section at 600°C

Fig. 13: Al-Ti-V.

Calculated section at

21 at.% V [1995Ahm]

52


MSIT®

Al–Ti–V

10 20 30 40

500

750

1000

1250

1500

1750

V 0.00

Ti 50.00

Al 50.00

V 50.00

Ti 50.00

Al 0.00V, at.%

Te

mp

era

ture

, °C

γ

α2+γ

α2+β0 α+β

β

β0

α

LFig. 14: Al-Ti-V.

Calculated section at

50 at.% Ti [1997Sha]

10

600

700

800

900

1000

V 0.00

Ti 94.80

Al 5.20

V 18.59

Ti 76.15

Al 5.26V, at.%

Te

mp

era

ture

, °C

α´ α´´ α´´+β β

Fig. 15: Al-Ti-V.

An isopleth at a

constant mass ratio of

Nb:Ti=1:4

53


MSIT®

Al–Ti–V

10

600

700

800

900

1000

1100

V 0.00

Ti 94.00

Al 6.00

V 20.00

Ti 74.00

Al 6.00V, at.%

Te

mp

era

ture

, °C

α´ α´´ α´´+β β

Fig. 16: Al-Ti-V.

Diagram of phase

compositions of

quenched Ti-6Al-V

(mass%) alloys

[1985Kol2]

80

90

10 20

10

20

Ti Ti 75.00

V 25.00

Al 0.00

Ti 75.00

V 0.00


Axes: at.%

2 1 1

1

α´

α´´

α´´+β+ω β+ω

β

α´´+β

Fig. 17: Al-Ti-V.

Projection of the

diagram of phase

compositions of

quenched Al-Ti-V

alloys [1985Kol2]

54


MSIT®

Al–Ti–Zr

Aluminium – Titanium – Zirconium

Ludmila Tretyachenko

Literature Data

Early studies of the Al-Ti-Zr phase diagram were carried out by X-ray diffraction (XRD) and differential

thermal analysis (DTA) in the Ti-Ti2Al-Zr region of the ternary system [1961San, 1964Kor1, 1964Kor2,

1964Pyl, 1968Shi1, 1968Shi2, 1970Nar, 1984Li]. The vertical sections Zr-(91Ti-9Al) (mass%), Zr-Ti3Al,

Zr-Ti2Al, (95Ti-5Zr)-Ti3Al (mass%) and the isopleth at 5 mass% Zr were presented and boundaries of the

+ 2 phase field at 500°C as well as the solubility limit of Al in the phase containing up to 10 at.% Zr at

600°C were determined. The crystal structure and hardening of the phase in the Ti-(48-54Al)-(0-12)Zr

(at.%) alloys were studied by [1987Kas]. Extensions and lattice parameters of the solid solution phases

(Ti1-xZrx)Al3 in the powdered as cast Al-2 at.% (Ti1-xZrx) (0 x 1) alloys were determined by [1982Tsu].

These works were reviewed by [1993Ans] and earlier by [1973Iva]. In addition, effect of 1 and 2 at.% Zr

addition on the /( + 2) phase boundary between 600 and 900°C was studied using optical microscopy

[1969Cro]. The influence of Zr additions up to 8 mass% on the /( + 2) phase boundary in the Ti-8 mass%

Al alloys containing oxygen (0.06 and 0.18 mass%) was studied by [1985Sca] using transmission electron

microscopy (TEM).

Following investigations of the Al-Ti-Zr system concerned Al rich alloys [1989Par, 1993Lee, 1997Fan,

2000Mal], alloys of the TiAl3-ZrAl3 section [2003Kar], TiAl based alloys [1988Has, 1992Che], alloys on

the base of Ti3Al [2000Sor] and phase equilibria between , ( 2), and phases [2000Kai].

Structure of the alloys containing 34, 36, and 38 mass% Al and up to 10 at.% Zr have been studied in as cast

and annealed at 1000°C states using optical microscopy, electron microprobe analysis (EMPA) and XRD

[1988Has]. Lattice parameters of the TiAl based phase ( ) and phase equilibria in the considered region

have been determined.

The isopleth at 2.3 at.% Zr and up to 10 at.% Al for the temperature range 800 - 1000°C has been presented

by [1988Gro]. Phase boundaries have been determined by EMPA and compared with those calculated from

the thermodynamic parameters.

Phase equilibria involving (hcp), (bcc) and (L10) phases have been studied in the temperature range

1000 - 1300°C and in the composition range of (0.5-12)Zr-Ti-(35-47)Al (at.%) [2000Kai]. The alloys were

prepared by arc melting and heat treated at 1000°C for 168 or 504 h, at 1200°C for 168 h and at 1300°C for

24 h followed by quenching into ice water. The study was made using optical microscopy and EMPA. The

appropriate partial isothermal sections at 1000, 1200 and 1300°C were determined.

Recently, the arc-melted alloys (ZrxTi1-x)Al3 (0 x 1) have been studied using in situ XRD in the

temperature range from 20 to 1100°C as well as by means of optical microscopy, DTA and Vickers hardness

tests [2003Kar].

The site occupancy of Zr in TiAl and Ti3Al based phases and phase stability of Ti2ZrAl have been studied

by means of the atom location channeling enhanced microanalysis (ALCHEMI) [1992Che, 1999Hao] and

by the first-principles electronic structure total energy calculations [1999Rav, 2002Rav]. The theoretical

and experimental investigations of sublattice substitution of alloying elements are summarized by

[2000Yan].

Binary Systems

The Al-Ti phase diagram is accepted from [1993Oka2], where the diagram was taken from the

thermodynamic assessment by [1992Kat]. This phase diagram was accepted also for the MSIT binary

evaluation program by [2003Sch2], where the TiAl-TiAl3 region determined by [2001Bra] was given in

addition.

The Al-Zr phase diagram is accepted from [1993Oka1]. The same Al-Zr phase diagram is given by

[2003Sch1].

The Ti-Zr phase diagram is taken from [Mas2].

55


MSIT®

Al–Ti–Zr

Solid Phases

The binary phases pertinent to the regions of the ternary Al-Ti-Zr system, which have been studied, are

given in Table 1. No ternary phases were found in the studied concentration and temperature ranges.

The hcp ( ) and bcc ( ) solid solutions exist in a wide range of compositions. All alloys in the Ti-Ti2Al-Zr

region at 1200°C have the bcc crystal structure.

A significant Zr solubility in Ti3Al ( 2) was observed. Although the maximum solubility has not been

determined, it is higher than 25 at.%, as confirmed by XRD of ZrTi2Al alloy, for which the D019 type crystal

structure has been found after annealing at 1000°C for 30 days [2000Sor].

The limit of Zr solubility in the phase (TiAl) is also not determined. It is more than ~11 at.% [2000Kai].

According to [1992Che], the 2Zr-50Ti-48Al (at.%) alloy produced by arc melting followed by

homogenization at 1100°C for 100 h, heat treatment at 800°C/3 h+600°C/3 h and air cooling to 25°C°C was

single phase .

The limited solid solutions on the base of TiAl3 and ZrAl3 compounds exist in the section TiAl3 - ZrAl3[1982Tsu, 2003Kar]. The D022 type structure (TiAl3) dissolves up to 2 at.% Zr, while the D023 type

structure (ZrAl3) exists in a wide range of compositions from pure ZrAl3 to about 15 at.% Ti [2003Kar]. Zr

and Ti substitute each other in both phases. Continuous solid solubility can be supposed between the

isostructural aluminides Ti5Al11 and ZrAl3, which both have the D023 type crystal structure. This follows

from the continuous variation of the lattice parameters of the D023 phase along the ZrAl3-TiAl3 section

almost over the whole concentration range.

The Ti solubility in other Zr aluminides has not been determined.

The metastable phase with the L12 (Cu3Au) type structure was obtained in the mechanical alloyed

ZrxTi25Al75-x (x up to 8) alloys [1997Fan]. Precipitates with the L12 metastable structure were found to

occur during aging the arc-melted Al alloy containing 1 vol.% (Zr0.75Ti0.25)Al3 [1989Par] and in rapidly

solidified Al-Ti-Zr alloys with 1.25 at.% (Ti+Zr) [2000Mal].

Liquidus Surface

The primary crystallization of the phase was observed over the whole Ti-Ti2Al-Zr region [1964Kor1,

1964Kor2, 1964Pyl]. The primary crystallization of the D023 phase with increasing melting temperature

from 1408 to 1607°C can be supposed over the TiAl3-ZrAl3 section [2003Kar].

Isothermal Sections

Partial isothermal sections at 1300 and 1200°C are shown in Figs. 1 and 2 from [2000Kai]. The isothermal

section at 1000°C (Fig. 3) is constructed using the data by [1988Has] for the 2+ region, [2000Sor]

concerning the homogeneity range of the 2 phase, as well as certain data by [1985Sca, 1988Gro] for the

Ti-rich region and the accepted Al-Ti phase diagram [1993Oka2, 2003Sch2]. The phase fields in the Ti

corner are shown tentatively because of shortage of reliable data.

The fragment of the isothermal section at 700°C shown by [1969Cro] evidences a decreasing Al solubility

in the phase with increasing Zr content in the alloys. The data by [1969Cro] are in good agreement with

the location of /( + 2) phase boundary at 700°C in the section Ti3Al-(25Ti-5Zr) (mass%) [1968Shi2] and

in the isopleth at 5 mass% Zr [1968Shi2, 1970Nar]. This boundary corresponds to ~12.5 at.% Al at ~1 at.%

Zr [1968Shi2, 1969Cro] and ~12 at.% Al at ~2 at.% Zr [1969Cro, 1968Shi2, 1970Nar]. The ( + 2)/ 2

phase boundary at 700°C corresponds to ~21 at.% Al at 0.5 at.% Zr [1968Shi2] and ~23.7 at.% Al at 2 at.%

Zr [1968Shi2, 1970Nar] that is in agreement with the accepted Al-Ti phase diagram.


Various temperature - composition sections presented by [1964Kor1, 1964Kor2, 1964Pyl, 1968Shi1,

1968Shi2, 1970Nar] as well as the /( + 2) phase boundaries in the isopleths at 0, 1 and 2 at.% Zr from

700 to 500°C by [1969Cro] are inconsistent with the Al-Ti system accepted at present, especially at

temperatures higher than 1000°C, and even with each other and, therefore, they are not reproduced here.

56


MSIT®

Al–Ti–Zr

The vertical section at 2.3 at.% Zr is drawn mainly from data of calculation [1988Gro] taking into account

the accepted Al-Ti phase diagram (Fig. 4).

The transformation temperature was found to exhibit the minimum at 660°C and ~65 mass% Zr

[1964Pyl] for the alloys in the ((91Ti-9Al) (mass%))-Zr section.


Relationships of hardness and resistivity versus composition have been studied for the alloys of the

Ti3Al-Zr, Ti2Al-Zr, (95Ti-5Zr (mass%))-Ti3Al, (91Ti-9Al (mass%))-Zr sections and the isopleth at 5

mass% Zr [1964Kor2, 1964Pyl, 1968Bor1, 1968Shi2, 1970Nar]. The highest hardness exhibited quenched

alloys with fine martensite-like structure. High temperature strength test has been carried out on the alloys

of the (91Ti-9Zr)-Ti3Al and with 5 mass% Zr sections [1968Shi2, 1970Nar]. Alloys with phase structure

have shown an increased creep rate. Fine dispersed 2 phase grains resulted in decreasing creep. However,

the alloys containing more than 10 mass% Al exhibited high brittleness. The alloy 5Zr-Ti-9Al (mass%) was

found to exhibit significant strength up to 700°C and favorable combination of strength and ductility at

room temperature. High temperature strength of the alloys in the (91Ti-9Al)-Zr (mass%) section has not

been decreased only up to 7 mass% Zr [1968Bor1].

Mechanical properties of the alloys containing up to 65 mass% Zr and 4, 6 and 7 mass% Al in the

temperature range from -196 to 700°C have been determined by [1968Bor2]. The alloys with 2 - 4 mass%

Al and 6 - 8 mass% Zr were shown to exhibit high ductility at low temperatures. The alloys with 6 mass%

Al and 20 mass% Zr can be used for long time at temperature below 500°C.

The influence of heat treatment on the properties of 2Zr-Ti-7Al (mass%) alloy has been studied by

[1975Mel].

Intensive oxidation of Ti2Al-Zr alloys in air was observed to start at temperatures 920 - 950°C at Zr content

15 - 45 mass%, 755°C at 50 mass% Zr and 600°C at 90 mass% Zr [1964Kor2]. A decrease of oxidation

resistance with increasing Zr content was also observed by [1968Bor2].

A study of electrical resistivity, temperature coefficients of resistivity and thermo emf in a couple with Cu

was carried out by [1976Kal].

Alloys of the Al-Ti-Zr system were found to be promising for development of dispersion strengthened

aluminium alloys, which can be applied up to 425°C, due to precipitations of the metastable L12 phase with

stable nanocrystalline microstructure [1989Par, 2000Mal].

An addition of Zr to Ti3Al was found to reduce the amount of absorbed hydrogen and to increase the 50 %

desorption temperature. The dehydrogenated alloys ZrxTi75-xAl25 (x = 15 and 25) have been turned into the

amorphous state [2001Ish].

Miscellaneous

The effect of Zr addition on lattice parameters of TiAl3 phase in mechanically alloyed 8Ti-Al (mass%) alloy

has been studied by XRD [1993Lee]. The 8Zr-25Ti-67Al (at.%) alloy produced by mechanical alloying has

been studied by means of XRD, TEM and differential scanning calorimetry (DSC) up to 700°C [1997Fan].

Rapidly solidified Al rich alloys with 1.25 at.% (Ti1-xZrx) (0 x 1) have been studied using EMPA,

optical microscopy, SEM, XRD and TEM [2000Mal]. A behavior of as cast 0.68Zr-0.3Ti-99.02Al (mass%)

((Zr0.75Ti0.25)Al3) alloy during aging has been studied using TEM [1989Par].

The as cast (ZrxTi1-x)Al3 alloys of the TiAl3-ZrAl3 section have been studied by [2003Kar] and the lattice

parameters of TiAl3 based phase (D022) were found to increase linearly in the concentration range of

0 x 0.08, that is in a good agreement with [1982Tsu], who has found a similar behavior of the lattice

parameters of D022 phase up to x = 0.11. The ZrAl3 based phase (D023) was found to be single phase in the

alloys with 0.4 x 1 and the lattice parameters decrease with decreasing x not only for single phase alloys

but for those in the two phase region D022+D023. The almost linear decreasing of D023 lattice parameters

was described with the empirical equations:

a = 399.8 - 7.7 (1 - x)

c = 1728 - 62.5 (1 - x)

57


MSIT®

Al–Ti–Zr

(a and c are given in pm). The linear decreasing of D023 lattice parameters was also found for 0.25 x 1

range [1982Tsu].

It was suggested that small additions of Si favor formation of the D022 structure [2003Kar]. Alloying of

TiAl3 with Zr and ZrAl3 with Ti was found to decrease hardness and increase ductility [2003Kar].

The alloys with 4 and 6 mass% Al and 20 - 65 mass% Zr at 600 and 700°C were found to be single phase

due to increasing transformation temperature by Al addition [1968Bor2].

Additions of Zr to TiAl3 were found to stabilize the metastable L12 structure obtained by mechanical

alloying [1997Fan]. The temperature of the L12 D022 phase transformation was measured to be ~422°C

for TiAl3. No phase transformations were observed in TiAl3 with the Zr additions up to 8 at.%

(ZrxTi25Al75-x), which have been annealed at temperatures up to 700°C.

Zr atoms occupy Ti sites for TiAl and Ti3Al alloys [1988Has, 1992Che, 1999Hao, 1999Rav, 2000Yan,

2002Rav].

Theoretical calculations of electronic structure and energy of three phases (B2, D019 and orthorhombic O)

for ZrTi2Al have shown ZrTi2Al is more stable in the D019 phase [1999Rav]. [2002Rav] using first

principles electronic structure total energy calculations to examine the phase stability of Ti2ZrAl has shown

that D019-like and L12-like structures of Ti2ZrAl are the competing ones among seven structures

considered.

The 2 (D019) structure was obtained for ZrTi2Al by [2000Sor], but ZrAl3 phase (L12) has been observed

in the alloys in the Ti2Al-Zr section annealed at 1100 and 500°C in a wide range of compositions by

[1964Kor2].

The influence of Zr on the lattice parameters of the phase is shown in Fig. 5 from [1988Has].

References

[1961San] Sandlin, D.R., Klung, H.A., Jr., “A Phase Study of a Selected Portion of the Ti-Al-Zr

Ternary System Including Lattice Parameter Determinations for the Ti-Al Phase”, Master

Thesis, Institute of Technology, Wright-Patterson Air Force Base, Ohio (1961) (Crys.

Structure, Experimental, 14) (quoted by [1993Ans])

[1964Kor1] Kornilov, I.I., Nartova, T.T., Savelïyeva, M.M., “Phase Equilibrium of Alloys in the

Ti3Al-Zr Section of the Ti-Al-Zr Ternary System” (in Russian), “Metallovedeniye Titana”,

Nauka, Moscow, 43-46 (1964) (Equi. Diagram, Experimental, #, 10)

[1964Kor2] Kornilov, I.I., Boriskina, N.G., “Study of the Phase Structure of Alloys of the Ti-Al-Zr

System along the Ti2Al-Zr Section” (in Russian), “Metallovedeniye Titana”, Nauka,

Moscow, 47-53 (1964) (Equi. Diagram, Experimental, 8)

[1964Pyl] Pylaeva, E.N., Volkova, M.A., “Investigation of Alloys of the Ternary Ti-Al-Zr System” (in

Russian), “Metallovedeniye Titana”, Nauka, Moscow, 38-42 (1964) (Equi. Diagram,

Experimental, 6)

[1968Bor1] Boriskina, N.G., Volkova, M.A., “Investigation of Alloys of the Ti-Al-Mo-Zr System by a

Bend Method at Elevated Temperature” (in Russian), “Titanovyye Splavy dlya Novoy

Tekhniki”, Nauka, Moscow, 164-171 (1968) (Experimental, 4)

[1968Bor2] Borisova, E.A., Shashenkova, I.I., “Investigation of Properties of Alloys of the Ti-Zr and

Ti-Zr-Al System” (in Russian), “Titanovyye Splavy dlya Novoy Tekhniki”, Nauka, Moscow,


[1968Shi1] Shirokova, N.I., Nartova, T.T., Kornilov, I.I., “Investigation of Equilibria and Properties of

Ti-Zr-Al Alloys” (in Russian), Izv. Akad. Nauk SSSR, Met., (4) 183-187 (1968) (Equi.

Diagram, Experimental, #, 15)

[1968Shi2] Shirokova, N.I., Nartova, T.T., “Investigation of Phase Equilibria and Properties of Alloys

of the Titanium Corner of the Ti-Zr-Al System” (in Russian), “Titanovyye Splavy dlya

Novoy Tekhniki”, Nauka, Moscow, 101-106 (1968) (Equi. Diagram, Experimental, 12)

[1969Cro] Crossley, F.A., “Effects of the Ternary Additions: O, Sn, Zr, Cb, Mo and V on the

/ +Ti3Al Boundary of Ti-Al Base Alloys”, Trans. Metall. Soc. AIME, 245, 1963-1968


58


MSIT®

Al–Ti–Zr

[1970Nar] Nartova, T.T., Shirokova, N.I., “Phase Equilibria and Heat Resistance of Al-Ti-Zr Alloys”

(in Russian), Izv. Akad. Nauk SSSR, Met., (3), 194-198 (1970) (Equi. Diagram,

Experimental, 12)

[1973Iva] Ivanov, O.S., Adamova, A.S., Tararayeva, E.M., Tregubov, I.A., (in Russian), “Structure of

Zr Alloys”, Nauka, Moscow, 56-57 (1973) (Equi. Diagram, Review, 4)

[1975Mel] Melnikova, V.I., Shklyar, R.Sh., Dyakonova, M.A., Potyomkina, T.G., Zvereva, Z.F.,

Kaganovich, I.N., “Influence of Composition and Heat Treatment on Properties of Alloys

of the Ti-Al System” (in Russian), Fiz. Met. Metalloved., 39, 1033-1036 (1975)

(Experimental, 7)

[1976Kal] Kalinin, G.P., Elyutin, O.P., Doronina, E.V., “Electrical Properties of Al-Ti-Zr Alloys” (in

Russian), Izv. Akad. Nauk SSSR, Met., (5), 220-223 (1976) (Experimental, 10)

[1981Kin] King, H.W., “Crystal Structure of the Elements at 25°C”, Bull. Alloy Phase Diagrams, 2,

401-402 (1981) (Crys. Structure, Review, 5)

[1982Tsu] Tsunekawa, S., Fine, M.E., “Lattice Parameters of Al3(ZrxTi1-x) vs. x in Al-2 at.% (Ti+Zr)

Alloys”, Scr. Metall., 16, 391-392 (1982) (Crys. Structure, Experimental, 2)

[1984Li] Li, D., Liu, Y.-Y., Wan, X.-J., “Thermal Stability of Titanium Alloys I. Electron

Concentration Rule for Formation ofthe Ti3X Phase” (in Chinese), Jinshu Xuebao, 20,

A375-A382 (1984) (Equi. Diagram, Experimental, 22)

[1985Sca] Scarr, G.K., Williams, J.C., Ankem, S., Bomberger, H.B., “The Effect of Zirconium and

Oxygen on -2 Precipitation in Titanium-Aluminum Alloys”, Titanium: Sci. Technol.,

Proc. Int. Conf. Titanium, 1984 (Pub. 1985), Luetjering, G., Zwicker, U., Bunk, W., (Eds.),

Dtsch. Ges. Metallkd., Oberursel, F.R.G., 3, 1475-1479 (1985) (Equi. Diagram,

Experimental, 3)

[1987Kas] Kasahara, K., Hashimoto, K., Doi, H., Tsujimoto, T., “Crystal Structure and Hardness of

TiAl Phase Containing Zr” (in Japanese), J. Jpn. Inst. Met., 51, 278-284 (1987) (Crys.


[1988Gro] Gros, J.P., Ansara, I., Allibert, M., “Prediction of / Equilibria in Titanium-Based Alloys

Containing Al, Mo, Zr, Cr. II”, Les Editions de Physique. Avenue du Hoggar, Zone

Industrielle de Courtaboeuf, B.P. 112, F-91944 Les Ulis Cedex, France, 6th World

Conference on Titanium. III, Cannes, France, 1559-1564 (1988) (Equi. Diagram,

Experimental, 0)

[1988Has] Hashimoto, K., Doi, H., Kasahara, K., Tsujimoto, T., Suzuki, T., “Effects of Third Elements

on the Structures of Ti-Al-Based Alloys” (in Japanese), J. Jpn. Inst. Met., 52, 816-825

(1988) (Crys. Structure, Equi. Diagram, Experimental, #, 31)

[1989Par] Parameswaran, V.R., Weertman, J.R., Fine, M.E., “Coarsening Behavior of L12 Phase in an

Al-Zr-Ti Alloy”, Scr. Metall., 23, 147-150 (1989) (Experimental, 9)

[1990Sch] Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”,

Z. Metallkd., 81, 389-396 (1990) (Crys. Structure, Equi. Diagram, Experimental, 33)

[1992Che] Chen, X.F., Reviere, R.D., Oliver, B.F., Brooks, C.R., “The Site Location of Zr Atoms

Dissolved in TiAl”, Scr. Metal. Mater., 27, 45-49 (1992) (Crys. Structure, Experimental, 5)


Ti-Al System”, Metall. Trans., 23A, 2715-2723 (1992) (Equi. Diagram, Review, Theory,

Thermodyn., 51)

[1993Ans] Ansara, I., Grieb, B., Legendre, B., “Aluminium - Titanium - Zirconium”, MSIT Ternary

Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science

International Services GmbH, Stuttgart; Document ID: 10.16126.1.20, (1993) (Crys.

Structure, Equi. Diagram, Assessment, 13)

[1993Lee] Lee, K.-M., Lee, J.-H., Moon, I.-H., “Effects of V and Zr Addition on Lattice Parameters of

Al3Ti Phase in Mechanically Alloyed Al-8 wt.% Ti Alloys”, Scr. Metall. Mater., 29,


[1993Oka1] Okamoto, H., “Al-Zr (Aluminium - Zirconium)”, J. Phase Equilib., 14, 259-260 (1993)

(Equi. Diagram, Review, #, 2)

59


MSIT®

Al–Ti–Zr

[1993Oka2] Okamoto, H., “Al-Ti (Aluminium - Titanium)”, J. Phase Equilib., 14, 120-121 (1993)

(Crys. Structure, Equi. Diagram, Review, #, 16)

[1996Tre] Tretyachenko, L.A., “On the Ti-Al System”, in “Phase Diagrams in Material Science”, 1th

International School, Katsyveli, Krimea, Ukraine, Abstracts, 118 (1996) (Equi. Diagram,

Experimental, #, 0)

[1997Bul] Bulanova, M., Tretyachenko, L., Golovkova, M., “Phase Equilibria in Ti-Rich Corner of the

Ti-Si-Al System”, Z. Metallkd., 88, 256-267 (1997) (Crys. Structure, Equi. Diagram,


[1997Fan] Fan, G.J., Song, X.P., Quan, M.X., Hu, Z.Q., “Mechanical Alloying and Thermal Stability

of Al67Ti25M8 (M = Cr, Zr, Cu)”, Mater. Sci. Eng., A231, 111-116 (1997) (Crys. Structure,

Experimental, 22)

[1998Hel] Hellwig, A., Palm, M., Inden, G., “Phase Equilibria in the Al-Nb-Ti System at High

Temperatures”, Intermetallics, 6, 79-84 (1998) (Crys. Structure, Equi. Diagram,

Experimental, 57)


Elements in TiAl and Ti3Al Alloys”, Acta Mater., 47, 1129-1139 (1999) (Crys. Structure,

Experimental, Theory, 41)

[1999Rav] Ravi, C., Vajeeston, P., Mathijaya, S., Asokamani, R., “Electronic Structure, Phase Stability

and Cohesive Properties of Ti2XAl (X = Nb, V, Zr)”, Phys. Rev. B, 60, 15683-15690 (1999)

(Crys. Structure, Theory, 32)

[2000Kai] Kainuma, R., Fujita, Y., Mitsui, H., Ohnuma, I., Ishida, K., “Phase Equilibria among

(hcp), (bcc) and (L10) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855-867

(2000) (Equi. Diagram, Experimental, #, 29)

[2000Mal] Milek, P., Janeoek, M., Smola, B., Bartuska, P., Plestil, J., “Structure and Properties of

Rapidly Solidified Al-Zr-Ti Alloys”, J. Mater. Sci., 35, 2625-2633 (2000) (Crys. Structure,

Experimental, 33)

[2000Sor] Sornadurai, D., Panigrahi, B., Sastry, V.S., Ramani, “Crystal Structure and X-Ray Powder

Diffraction Pattern for Ti2AlZr”, Powder Diffr., 15, 189-190 (2000) (Crys. Structure,

Experimental, 6)

[2000Yan] Yang, R., Hao, Y., Song, Y., Guo, Z.X., “Site Occupancy of Alloying Additions in Titanium

Aluminides and Its Application to Phase Equilibrium Evaluation”, Z. Metallkd., 91,

296-301 (2000) (Crys. Structure, Theory, Review, 38)


Binary System Ti-Al”, Metall. Mater. Trans., 32A, 1037-1047 (2001) (Crys. Structure,


[2001Ish] Ishikawa, K., Hashi, K., Suzuki, K., Aoki, K., “Effect of Substitutional Elements on the

Hydrogen Absorption-Desorption Properties of Ti3Al Compounds”, J. Alloys Compd., 314,


[2002Rav] Ravi, C., Mathijaya, S., Valsakumar, M.C., Asokamani, R., “Site Preference of Zr in Ti3Al

and Phase Stability of Ti2ZrAl”, Phys. Rev. B, 65, (155118-1)-(155118-6) (2002) (Crys.

Structure, Theory, 34)

[2003Kar] Karpets, M.V., Milman, Yu.V., Barabash, O.M., Korzhova, N.P., Senkov, O.N., Miracle,

D.B., Legkaya, T.N., Voskoboynik, I.V., “The Influence of Zr Alloying on the Structure and

Properties of Al3Ti”, Intermetallics, 11, 241-249 (2003) (Crys. Structure, Equi. Diagram,

Experimental, 16)

[2003Sch1] Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Service, GmbH,

Stuttgart, to be published, (2003) (Equi. Diagram, Review, 85)

[2003Sch2] Schuster, J.C., “Al-Zr (Aluminium-Zirconium)”, MSIT Binary Evaluation Program, in


GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram,

Assessment, 151)

60


MSIT®

Al–Ti–Zr


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Al)

< 660.452

cF4

Fm3m

Cu

a = 404.96

a = 404.9 0.4

a = 405.1 0.4

a = 404.2 0.4

0.6 at.% Ti at 664.2°C [1993Oka2] and

0.08 at.% Zr at 660.8°C [1993Oka1]

pure Al at 25°C [1981Kin, V-C2]

in the melt-spun ribbon of 4Zr-Al

(mass%) [2000Mal]

in the melt-spun ribbon of

2.1Zr-1.1Ti-Al (mass%) [2000Mal]

in the melt-spun ribbon of 2.2Ti-Al

(mass%) [2000Mal]

, Ti1-xZrx(h)

Ti(h)

1670 - 882

Zr(h)

1885 - 863

cI2

Im3m

W

a = 330.65

a = 360.9

0 x 1 [Mas2, V-C2]

dissolves up to 44.8 at.% Al at x = 0 and

1490°C [1993Oka2] and up to 25.9 at.%

Al at x = 1 and 1350°C [1993Oka1]

[Mas2]

[Mas2]

, Ti1-xZrx(r)

Ti(r)

< 882

Zr(r)

< 863

hP2

P63/mmc

Mg

a = 295.03

c = 468.36

a = 323.17

c = 514.76

[Mas2, V-C2]

47.3 to 51.4 at.% Al at x = 0 at 1463°C

[1993Oka2] and 0 to 8.3 at.% Al at x = 1

at 910°C [1993Oka1]

pure Ti at 25°C [Mas2, 1981Kin]

pure Zr at 25°C [Mas2, 1981Kin]

2, (Ti1-xZrx)3Al

Ti3Al

1164

hP8

P63/mmc

Ni3Sn

a = 580.6

c = 465.5

a = 580.6

c = 465.5

a = 596.1 0.1

c = 479.3 0.1

(D019) ordered phase [V-C]

0 to 38.2 at.% Al, maximum at 30.9 at.%

and 1164°C at x = 0 [1993Oka2,

2003Sch2]

at 22 at.% Al [L-B]

at 38 at.% Al [L-B]

single phase ZrTi2Al annealed at

1000°C for 30 d [2000Sor]

61


MSIT®

Al–Ti–Zr

, (Ti1-xZrx)1-yAly

TiAl

< 1463

tP4

P4/mmm

CuAu

a = 400.5

c = 407.0

a = 400.0 0.1

c = 407.5 0.1

a = 398.4 0.1

c = 406.0 0.1

L10 ordered phase

46.7 to 66.5 at.% Al [1993Oka2,

2003Sch2]

0 to 62 at.% Al at 1200°C [2001Bra]

[V-C]

at 50 at.% Al [2001Bra]

at 62 at.% Al [2001Bra]

, TiAl2< 1199

oC12

Cmmm

ZrGa2

tP4

P4/mmm

CuAu

tI24

I41/amd

HfGa2

tP32

P4/mbm

Ti3Al5

a = 1208.84

b = 394.61

c = 402.95

a = 403.0

c = 395.5

a = 397.0

c = 2430.9

a = 396.7

c = 2429.68

a = 1129.3

c = 403.8

chosen stoichiometry [1992Kat,

1993Oka2], summarizes several phases

[2003Sch2]:

metastable modification of TiAl2,

observed only in as cast alloys

[2001Bra], listed as TiAl2(h) (66 to 67

at.% Al, 1433 - 1214°C) by [1990Sch]

Ti1-xAl1+x, 63 to 65 at.% Al at 1300°C,

stable in the range 1445 - 1170°C

[2001Bra], listed orthorhombic, Pmmm,

with pseudotetragonal cell, stable in the

range ~1445 - 1424º by [1990Sch]

for Ti36Al34 at 1300°C [2001Bra]

stable structure of TiAl2< 1215°C [2001Bra],

listed as TiAl2(r) by [1990Sch]

[2001Bra]

[1990Sch]

Ti3Al5, stable below 810°C [2001Bra]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

62


MSIT®

Al–Ti–Zr

2

tetragonal

superstructure of

CuAu type

tI16

I4/mmm

ZrAl3

tP28

P4/mmm

Ti2Al5

a* = 395.3

c* = 410.4

a* = 391.8

c* = 415.4

a = 393.81 to 392.3

c = 1649.69 to

1653.49

a = 393

c = 1654

a = 390.53

c = 2919.63


Ti5Al11 [2001Bra] stable in the range

1416 - 995°C, 66 to 71 at.% Al at

1300°C (including the stoichiometry

Ti2Al5) [2001Bra]

at 66 at.% Al [2001Bra]

* CuAu subcell only

at 71 at.% Al [2001Bra]

* CuAu subcell only

D023 type [V-C]

68.5 to 70.9 at.% Al, 1416 - 1206°C

[1990Sch]

69 - 71 at.% Al, 1450 - 990°C [1996Tre,

1997Bul]

for Ti-69.4Al (at.%), accepted as Ti2Al5,

stable between 69.4 and 71.8 at.% Al at

1200°C [1998Hel]

“Ti2Al5”, 1416 - 990°C [1992Kat,

1993Oka2, 2003Sch2] 1216 - 985°C

[1990Sch] included in the homogeneity

range of Ti5Al11 [2001Bra]

, (Ti1-xZrx)Al3

TiAl3(h)

< 1393

tI8

I4/mmm

TiAl3

a = 384.9

c = 860.9

a = 385.3 0.2

c = 861.8 0.2

a = 385.8

c = 858.7

a = 385.3 to 386.2

c = 858.7 to 865.0

a = 386.3 0.2

c = 864.8 0.3

a = 389.3 0.3

c = 871.4 0.5

a = 389.8 0.4

c = 872.6 0.6

(D022) [V-C]

0 x 0.11 (0 to ~2.75 at.% Zr) for as

cast alloys Al-2 at.% (Ti1-xZrx) 0 x 1

[1982Tsu], [1992Kat, 1993Oka2]

from 1387 to ~950°C for the Ti rich

region, from 1387 to 735°C for the Al

rich region, homogeneity range 74.5 to

75 at.% Al at 1200°C [2001Bra]

melting temperature 1408°C [2003Kar],

< 1425°C [1996Tre]

[2001Bra]

in as cast Al+2 at.% Ti and in Ar

atomized 4.7Ti-Al (at.%) alloys

[1982Tsu]

in mechanically alloyed Al-8 mass%

(4.7 at.%) Ti [1993Lee]

in the as cast alloys at 0 x 0.08

[2003Kar]

in the (Ti0.84Zr0.16)Al3 alloy at 20°C

[2003Kar]


[2003Kar]


[2003Kar]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

63


MSIT®

Al–Ti–Zr

TiAl3(l)

< 950 (Ti rich)

< 735 (Al rich)

tI32

I4/mmm

TiAl3(l)

a = 387.7

c = 3382.8

a = 387.5

c = 3384

74.5 to 75 at.% Al [2001Bra]

listed as Ti8Al24, exists < 650°C

[2003Kar]

(Ti1-xZrx)Al3

TiAl3

cP4

Pm3m

Cu3Au

a = 397.2

a = 399 1

a = 404 1

a = 405 1

a = 408 1

a = 403.3

metastable phase L12 in Al rich alloys

[1997Fan, 1989Par, 2000Mal]

obtained from splat cooling for Ti-85Al

(at.%) [2001Bra]

at x = 0.25, as melt-spun ribbon of

1.0Zr-1.65Ti-Al (mass%) [2000Mal]


2.1Zr-1.1Ti-Al (mass%) [2000Mal]


3.1Zr-0.55Ti-Al (mass%) [2000Mal]

at x = 1, as melt-spun ribbon of

4.1Zr-Al (mass%) [2000Mal]

in Zr8Ti25Al67 alloy after milling for

40 h [1997Fan]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

64


MSIT®

Al–Ti–Zr

1, (TixZr1-x)Al3

ZrAl3 < 1580

tI16

I4/mmm

ZrAl3

a = 401.4

c = 1732

a = 399.8 0.1

c = 1727.6 0.3

a = 397.8 0.1

c = 1718.5 0.4

a = 396.3 0.1

c =1708.4 0.3

a = 396.0 0.1

c = 1690.7 0.3

a = 394.3 0.1

c = 1684.7 0.4

a = 392.9 0.4

c = 1679.2 0.5

a = 393.7

c = 1677.0

a = 393.3 0.2

c = 1673.2 0.3

a = 393.2 0.2

c = 1670.7 0.6

a = 392.7 0.3

c = 1668.4 0.5

a = 394.9 0.2

c = 1693.1 0.4

a = 395.3 0.3

c = 1695.0 0.6

D023 type, 0 x 0.75 for as cast alloys

Al-2 at.% (TixZr1-x) [1982Tsu]

single phase in as cast (TixZr1-x)Al3alloys at 0 x 0.4 [2003Kar]

[1993Oka1, V-C]

melting temperature 1607°C [2003Kar]

x = 0.2 [2003Kar]

x = 0.4 [2003Kar]

in the as cast (Ti0.6Zr0.4)Al3 alloy

(D023+D022) [2003Kar]


(D023+D022) [2003Kar]


(D023+D022) [2003Kar]

in the mechanically alloyed

Al-8Ti-3.8Zr (mass%) ((Ti0.8Zr0.2)Al3)

[1993Lee]


(D022+D023) [2003Kar]


(D022+D023) [2003Kar]


(D022+D023) [2003Kar]


(D023+D022) [2003Kar]

in the same alloy (D023+D022) at 700°C

[2003Kar]

Zr3Al

< 1019

cP4

Pm3m

Cu3Au

a = 437.2 (L12) [1993Oka1, V-C, 2003Sch1]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

65


MSIT®

Al–Ti–Zr

10

20

30

50 60 70

30

40

50

Zr 40.00

Ti 40.00

Al 20.00

Zr 0.00

Ti 80.00

Al 20.00

Zr 0.00

Ti 40.00


Axes: at.%

γ

α

β

γ+βγ+α

Fig. 1: Al-Ti-Zr.

Partial isothermal

section at 1300°C

10

20

30

50 60 70

30

40

50

Zr 40.00

Ti 40.00

Al 20.00

Zr 0.00

Ti 80.00

Al 20.00

Zr 0.00

Ti 40.00


Axes: at.%

γ

β

α

γ+βγ+α

Fig. 2: Al-Ti-Zr.

Partial isothermal

section at 1200°C

66


MSIT®

Al–Ti–Zr

10

20

30

40

50

50 60 70 80 90

10

20

30

40

50

Zr 60.00

Ti 40.00

Al 0.00

Ti

Zr 0.00

Ti 40.00


Axes: at.%

γ

α2

β

α

Fig. 3: Al-Ti-Zr.

Partial isothermal

section at 1000°C

10

700

800

900

1000

Al 0.00

Ti 97.70

Zr 2.30

Al 20.00

Ti 77.70

Zr 2.30Al, at.%

Te

mp

era

ture

, °C

αα+α2

β

Fig. 4: Al-Ti-Zr.

Partial section at 2.3

at.% Zr

67


MSIT®

Al–Ti–Zr

Zr, at.%

La

ttic

epa

ram

ete

r,p

m

4

400

Zr 0.00

Al 48.00

Ti 52.00

Zr 12.00

Al 42.00

Ti 46.00

8

402

404

406

408

410

c

a

Fig. 5: Al-Ti-Zr.

Influence of Zr on the

lattice parameters of

the phase in the

(52Ti-48Al)-Zr

section [1988Has]

68


MSIT®

B–C–Ti

Boron – Carbon – Titanium

Peter Rogl, Hans Bittermann and Hans Duschanek

Literature Data

Due to chemical compatibility with titanium metal, both titanium boride and titanium carbide have become

attractive discontinuous reinforcements in titanium matrix composites for high-temperature structural

applications. Therefore early studies were devoted to elucidate the interaction between titanium boride,

titanium carbide and carbon or boron carbide [1951Gre, 1952Gla, 1955Bre, 1956Gea, 1959Sam, 1960Por].

Based on these findings the authors of [1961Now] established phase relations in the B-C-Ti ternary system

from X-ray powder analyses of about 50 alloys hot pressed from the elements in graphite dies and annealed

at 1500, 1800 or 2000°C, respectively. Preliminary information on the pseudobinary eutectic systems

TiB2-TiC1-x, TiB2-‘B4C’ and TiB2-C were provided by [1956Gea, 1959Sam, 1960Por, 1965Lev].

The most complete experimental information is from [1966Rud], comprising a reinvestigation of the

isothermal section at 1400°C, a determination of the liquidus surface and an investigation of three isopleths,

TiB2-TiC0.92, TiB2-C and TiB2-B4.5C. On the basis of these results [1966Rud] derived a reaction scheme

for the ternary system and constructed a series of further isotherms at 1500, 1600, 1700, 2000, 2160, 2300,

2420, 2600 and 2800°C as well as a three-dimensional isometric view of the ternary system. A shortcoming

of the data presented by [1966Rud] is the missing information on the binary stable phase Ti3B4 described

by [1966Fen].

For their investigations [1966Rud] employed X-ray powder diffraction, Pirani melting point,

metallographic and differential thermoanalytical (> 4 K s-1) techniques on hot pressed and sintered as well

as argon arc melted specimens. Starting materials were high purity elemental powders (i.e. Ti containing

1300 ppm C, 50 ppm N, 200 ppm O; spectrographic grade graphite powder with impurities less than 100

ppm and boron powder of 99.55 mass% B containing 0.25 mass% Fe and 0.1 mass% C) as well as

prereacted master alloys of TiB2 (65.3 at.% B with 0.088 mass% C) and TiC1-x (powder with a particle size

< 80 m containing 49.4 at.% C of which 0.5 at.% was in the form of elemental carbide, a = 432.3 pm).

Pirani and DTA measurements were calibrated against an internal laboratory standard of the W-W2C

eutectic at 2710 20°C [1965Rud]. A total of 200 alloys were prepared by [1969Rud] mainly by short

duration hot pressing in graphite dies at temperatures between 1800 and 2200°C. After removing the surface

reaction zones, the samples were directly used in as-pressed condition for Pirani melting point (under 2.5

105 Pa He) or differential thermal analysis (in graphite container under 105 Pa He). Selected alloys from the

metal-rich region (> 85 at.% Ti) intended for melting point or DTA studies were electron or arc melted prior

to the runs. Whereas specimens for DTA and melting point analyses were directly equilibrated in the

equipment prior to the runs, specimens for the isothermal sections were generally annealed in a tungsten

mesh furnace for 68 h at 1400°C under a vacuum of 2 10-3 Pa or those from the concentration range

TiB2-TiC1-x-B-C for 12 h at 1750°C under 1.05 to 2 105 Pa of helium and rapidly quenched. Some alloys

were equilibrated in the melting point furnace and quenched in a preheated tin bath (300°C). Samples were

chemically analyzed for free and bound carbon, boron as well as oxygen and nitrogen contaminants. For

polishing and etching usually a slurry of alumina in 5% chromic acid was used; for alloys within the nominal

composition Ti-Ti0.7C0.3-Ti0.5C0.2B0.3-Ti0.7B0.3 anodic oxidation in an electroetching process using 10%

oxalic acid was said to provide excellent phase contrast coloring the metal phase light blue, the monoboride

brown, whereas the diboride remained unaffected. Specimens from the range

Ti0.7C0.3-Ti0.5B0.3C0.2-B0.3C0.7 were dip-etched in a solution with 1% aqua regia and HF. Samples from

the region Ti0.7B0.3-Ti0.2B0.3C0.5-Ti0.2B0.5C0.3-Ti0.5B0.5 were prepolished and etched with Murakamis

etchant, followed by dip-etching in 1% aqua regia + HF [1966Rud].

Independent studies of the isopleths TiB2-TiC1-x [1964Ord, 1975Ord] and TiB2-’B4C’ [1986Ord]

confirmed the pseudobinary eutectic nature of these sections. However, eutectic compositions and eutectic

temperatures shown by [1975Ord, 1986Ord] are at considerable variance to the findings of [1966Rud]. The

low eutectic temperatures reported by [1975Ord, 1986Ord] are likely due to insufficient correction for

69


MSIT®

B–C–Ti

non-black body conditions. A similar explanation may also hold for the low eutectic temperature

(TE = 2290°C) recorded by [1965Lev] for TiB2-C, for which the value 2507 15°C derived by [1966Rud]

seems to be more accurate. Solid solubility limits of TiB2 in TiC1-x have been studied by [1975Aiv] using

samples crystallized from the gas phase. Whereas a schematically correct isothermal section at 2000°C

including binary Ti3B4 was presented by [1985Cor], their isopleth Ti-’B4C’ was still based on the data by

[1966Rud] disregarding the interaction and phase relations with Ti3B4.

A preliminary critical assessment of the B-C-Ti system due to [1992Ale] refers to an outdated B-C binary

system involving phases such as B13C2 and B12C3 as individual compounds rather than compositions within

the homogeneous range of ‘B4C’. Compilations of the B-C-Ti system were presented by [1972Upa,

1983Sco, 1984Hol] and [1995Vil]. A thermodynamic calculation of the phase relations in the B-C-Ti

system was attempted by [1990Spe] in order to predict metastable phase formation during PVD.

Calculations for two isothermal sections at 2000 and 2300°C were presented as well as Gibbs energy curves

at 2270°C for the sections TiB2-TiC1-x and TiB2-C. Assuming sub-regular solutions, a thermodynamic

estimation of the complete B-C-Ti ternary was reported by [1997Gus] for the region from 27 to 3227°C. A

first thermodynamic calculation of the complete B-C-Ti ternary is due to [1995Dus]. With respect to recent

thermodynamic assessments of the binary systems B-C [1996Kas] and C-Ti [1997Dum], a new modelling

of the ternary was presented [1998Bit] and follows the outline given by [1995Dus].

Binary Systems

The B-Ti phase diagram first established by [1966Rud] is accepted from a recent assessment by [1986Mur]

including the boride Ti3B4 not reported in [1966Rud]; the most recent thermodynamic calculation was

performed by [1994Bae]. The phase diagram of the C-Ti system is essentially based on [Mas2]; the high

temperature phase ‘Ti2C’, however, has been omitted as there are severe doubts about its existence free

from non-metal contaminants. The low temperature ordering phases ‘Ti2C’, ‘Ti6C5’, ‘Ti3C2’, were not

included; their existence was proposed by [1991Gus] from theoretical thermodynamic estimations.

Thermodynamic descriptions for the C-Ti and B-C system are due to [1997Dum] and [1996Kas],

respectively. The crystallographic data for the binary boundary phases including solid solutions extending

into the ternary are presented in Table 1. It should be noted that the compositions of the liquidus l/l+TiB2

and l/l+TiC1-x, as reported by the various research groups, reveal a remarkable scatter or even less likely

positive derivatives d2T/dx2 due to the lack of suitable crucible materials for the aggressive Ti- or

B- containing melts at elevated temperatures. For consistency with the reported ternary, this assessment will

follow the experimental version presented by [1966Rud]. However, one has to keep in mind that the Pirani

melting point technique used by [1966Rud] and others is virtually incapable of providing reliable liquidus

data because of the gradual loss of mechanical strength of the sample specimen during measurement and

due to the loss of the black body measuring-hole.

Solid Phases

No ternary compounds exist in the B-C-Ti ternary system. Mutual solid solubility among the binary boride

and carbide phases generally was found to be very small [1961Now, 1965Lev, 1966Rud, 1975Ord,

1986Ord] except for the titanium monocarbide, for which lattice parameters in the ternary are considerably

increased with respect to those of the binary [1966Rud]. The large decrease of lattice parameters in TiB2 (in

B4C samples with 0 to 50 mass% TiB2 at 2160°C) was attributed by the authors to the incorporation of W

impurities from the WC-lined ball milling system [1985Nis]. A maximal solubility of 9 mass% TiB2 in B4C

at 2150°C was reported [1985Nis]. The ternary solid solubility of the non-metal elements in ( Ti) at the

temperature of the ternary eutectic L ( Ti)+TiB+TiC1-x was said to be smaller than 1 at.% B and 1 at.% C,

respectively [1966Rud]. The formation of binary Ti3B4 was confirmed from binary and ternary alloys

annealed in the temperature region 1400 to 1800°C [1991Pak], at 1550°C [1995Dus] and at 1600°C

[1996Bro]; lattice parameters in binary and ternary alloys compare well suggesting a very limited,

practically negligible solubility of C in Ti3B4 [1995Dus]. Single phase deposits if boron-containing titanium

carbide with boron contents as high as TiC0.81B0.17 were obtained via crystallization from the gas phase

(TiCl4, CCl4, BCl3) in the temperature range from 1100 to 1700°C on a molybdenum substrate [1975Aiv].

70


MSIT®

B–C–Ti

After prolonged annealing at 1200°C the X-ray diffractograms revealed TiB2 as a secondary phase

[1975Aiv].


A reaction scheme for the ternary B-C-Ti system including ten ternary invariant equilibria was provided by

[1966Rud]. Figure 1 shows our calculated version. Table 2 lists the compositions of the phases at the

four-phase isothermal reactions as given by [1966Rud] and compares experimental data with the results of

the thermodynamic calculation including also the reactions involving Ti3B4. As far as the very boron-rich

region is concerned the thermodynamic calculation favors a transition type reaction L+‘B4C’ TiB2+(B) at

2085°C rather than a ternary eutectic L (B)+‘B4C’+TiB2 at 2016°C as reported by [1966Rud]. This

discrepancy essentially results from the assessment of the B-C system by [1996Kas] revealing a peritectic

reaction L+‘B4C’ (B) at 2103°C rather than a eutectic L (B)+‘B4C’ at 2080°C as given by [1966Rud].

Also in the titanium-rich region the thermodynamic calculation yields a ternary peritectoid reaction:

( Ti)+TiB+TiC1-x ( Ti), at 920°C, instead of the transition reaction: ( Ti)+TiC1-x TiB+( Ti) at 890°C,

assumed by [1966Rud]. A DTA-run [1966Rud] performed on the ternary alloy Ti80B10C10 indicated a

( Ti)/( Ti)-transformation temperature 880 Ttr 930°C, which lay between those of the corresponding

reaction isotherms in the binary boundary systems ( Ti)+TiB ( Ti): [1966Rud] 880°C, [1994Bae] 883°C;

( Ti)+TiC1-x ( Ti): [1966Rud] 930°C, [1997Dum] 919°C). Whereas the thermodynamic calculation of

the B-Ti system closely adheres to the experimental value, the modeling of the C-Ti binary yields a

considerable reduction of the peritectoid temperature for the formation of ( Ti) from ( Ti)+TiC1-x, thus

resulting in a ternary peritectoid formation of ( Ti) at 920°C (see Table 2).

Liquidus Surface

Figure 2 is a representation of the liquidus surface based on results of [1966Rud] with small changes

referring to the existence of Ti3B4. It should be mentioned that the isotherms in the liquidus projections near

the B-Ti and the C-Ti boundary systems are dependent on the selected slope of the l/l+TiB2 and l/l+TiC1-x

liquidus (see section “Binary Systems” and compare with calculated liquidus projection shown in Fig. 3).

Isothermal Sections

Based on the experimentally determined isothermal section at 1400°C, as well as on the liquidus projection

derived (Fig. 2) and from the phase relations experimentally established for three isopleths, TiB2-C (Fig. 4),

TiB2-TiC0.92 (Fig. 5) and TiB2-B4.5C (Fig. 6), [1966Rud] constructed a series of isothermal sections at

1500, 1600, 1700, 2000, 2160, 2300, 2420, 2600 and 2800°C. In order to comply with the experiments of

[1985Cor, 1991Pak, 1995Dus, 1996Bro] minor corrections are needed to include the existence of Ti3B4; the

schematic representation of the phase relations at about 2000°C [1985Cor] basically agrees with the data of

[1966Rud]. [1995Dus] and [1996Bro] determined the exact position of the narrow two-phase field,

Ti3B4+TiC0.65, at 1550 and 1600°C, respectively. For comparison with the thermodynamic calculations see

section “Thermodynamics”.


Three pseudobinary systems of the eutectic type, e2, e3, e5, were established by [1966Rud]; they are

presented in Figs. 4, 5, 6, (compare also Table 2 and Fig. 1). The eutectic nature of the TiB2-C pseudobinary

was earlier reported by [1965Lev], the eutectic temperature, 2290 30°C, as measured by optical pyrometry

on prereacted powders through a bore hole in a directly heated graphite tube, is remarkably low compared

to 2507 15°C obtained by [1966Rud]. Probably the insufficient correction for non-black body conditions

in the experiment of [1965Lev] explains his low eutectic temperature. The eutectic composition in the

isopleth TiB2-C at ~85 mol% C [1965Lev] (originally given at ~85 at.% C; the conversion from at.% to

mol % would yield 94.5 mol% C) is in distinct disagreement with the eutectic at 58.6 mol% C (32 2 at.% C)

[1966Rud], which is accepted in this assessment. The compositional data in all isopleths given by

71


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B–C–Ti

[1966Rud] refer to 1 mole of atoms alloy in contrast to the chemical formula shown in the original figures

of the authors (i.e. label TiB2 should read as Ti0.33B0.67).

Recrystallization of the TiB2-TiC1-x pseudobinary eutectic at temperatures close to the eutectic line was said

to be fast, and fully or partially annealed structures resulted if cooling rates lower than 50 K s-1 were

employed [1966Rud]. Alloys along the section TiB2-TiC1-x (0 < x < 0.28) after annealing at 1400°C were

found to be single phase, if they contained less than 2 at.% B. The monocarbide was said to dissolve a

maximum of 5 mol% TiB2 at the temperature of the eutectic TiB2+TiC0.92 (TE = 2620 15°C), and solubility

is larger in the carbon-deficient carbide with boron atoms substituting for carbon and additionally filling

vacancies. TiB2 precipitation from the boron containing titanium-carbide was observed to be much faster

from the carbon-richer compositions, and cooling rates faster than 40 K s-1 were required to prevent

dissolution reactions [1966Rud]. The solubility of C in the diboride phase was said to be less than 2 mol%

C in 1 mole of atoms TiB2 [1966Rud]. Both the eutectic composition and the eutectic temperature (43.6

mol% TiC0.95 at 2520 40°C) reported by [1975Ord] are in rather poor agreement with the data recorded by

[1966Rud] for the pseudobinary eutectic TiB2-TiC0.92 (67.5 2 mol% TiC0.92 at 2620 15°C, originally

given as 57 2 mol% Ti0.52C0.48 by the authors). The lower temperature of [1975Ord] may eventually be

due to insufficient correction for non-black body conditions in pyrometric recording. Acceptable agreement

exists on the maximal solid solubility of 3.3 mol% TiB2 in TiC0.95 at 2620°C (originally given as 5 mol%

Ti0.33B0.67 [1966Rud]) and 2.6 to 3.4 mol% TiB2 in TiC0.95 at 2520°C by [1975Ord]. The eutectic

temperatures were said to fall for TiC1-x with x from 2520°C (x = 0.05) to 2380°C (x = 0.32), whereby the

maximal solid solubility of TiB2 in TiC1-x increases from ~3.5 mol% TiB2 for x = 0.05 to ~7 mol% TiB2

for x = 0.32 [1975Ord] (for lattice parameters see Table 1). However, it has to be noted, that both studies

[1966Rud, 1975Ord] did not choose the correct maximum melting point of TiC1-x at Ti55C45 TiC0.82

(calculated by [1997Dum] at Ti0.56C0.44 TiC0.79) required for a true pseudobinary. A directionally

solidified eutectic structure TiC-TiB2 was reported by [1980Ber].

Agreement exists on the eutectic nature of the TiB2-’B4C’ section investigated by [1966Rud]

(TE = 2310 15°C at 80 3 mol% B4.5C, originally given as 88 mol% B0.817C0.183) and [1986Ord]

(TE = 2200 40°C at 78 mol% ‘B4C’). However, correspondence of the eutectic parameters is poor. The data

reported by [1960Por], i.e. TE 1900°C for 75 mol% TiB2 have to be taken with caution.

Thermodynamics

The thermodynamic calculation [1998Bit] of the ternary B-C-Ti system by means of the THERMOCALC

program relied on the thermodynamic assessments of the binary systems B-C [1996Kas], B-Ti [1994Bae]

and C-Ti [1997Dum]. The liquid phase was described by adopting a substitutional solution model

[1990Sun] with a single sublattice (B,C,Ti)1. Titanium carbide was treated by [1997Dum] as an interstitial

solid solution of carbon in fcc-Ti. According to the experimental data of [1966Rud] the solubility of B in

TiC1-x is not negligible. Thus the sublattice model of TiC1-x was extended to a mixture of carbon atoms,

boron atoms and vacancies in the non-metal sites. Both the Ti (hcp) and Ti (bcc) phases are interstitial solid

solutions modeled with a two-sublattice model. The first sublattice is filled with titanium and on the second

sublattice, which represents the interstitial sites in the ternary, boron, carbon and vacancies are mixing. B

was treated by [1996Kas] as (B)93(B,C)12. Graphite was considered by [1996Kas] as a substitutional phase

with a single sublattice: (C,B)1. The ‘B4C’ phase was modeled by [1996Kas] with two sublattices, the

icosahedral lattice filled with the species B12 and B11C, and the other with the species B2, C2B and B2C:

(B12,B11C)1 (B2,C2B,B2C)1. No ternary solubility of Ti in ‘B4C’ was assumed. The borides TiB, Ti3B4 and

TiB2 were treated by [1994Bae] as stoichiometric phases and were described with a two sublattice model.

The first sublattice is completely filled with Ti atoms, the second one with B atoms: (Ti)1(B)1, (Ti)3(B)4

and (Ti)1(B)2. Only phase diagram data from [1966Rud] were selected for the optimization of the

thermodynamic parameters. The calculated liquidus surface is shown in Fig. 3. Various calculated

isothermal sections are shown in Figs. 7, 8, 9, 10, vertical sections in Figs. 11, 12, 13 and 14. Note: the

isopleths Ti-B0.5C0.5, Ti0.5C0.5-B and Ti0.5B0.5-C were constructed by [1966Rud] consistent with the

experimental data derived from the isothermal sections at 1400°C, the liquidus projection and the

72


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B–C–Ti

concentration sections TiB2-C, TiB2-TiC1-x and TiB2-’B4C’. These isopleths in the original version lack

information concerning the phase Ti3B4.


There are numerous papers and patents dealing with the technical application of TiC1-x/TiB2/’B4C’

products. We wish to emphasize, that in this article we essentially focus on the phase relations of the B-C-Ti

ternary only.

Both TiC and TiB2 are promising hard layer materials for wear protection. Abrasive mechanical properties

from samples TiC1-x-TiB2 have been reported by [1982Unr]. Indentation hardness, toughness and wear

resistance were observed as functions of the interlamellar spacing of directionally solidified eutectic

compositions TiB2/TiC1-x [1980Stu]. A series of papers deals with the increase of wear behavior of (a)

magnetron sputtered TiB2/TiC1-x coatings with a composition near the ternary TiC-TiB2 eutectic on

cemented [1991Hol], (b) codeposited TiB2/TiC1-x coatings on Ta-substrates from the gas phase (TiCl4,

n-C7H16, BCl3, H) in a cold wall reactor [1991Bar, 1995Gui], c) magnetron sputtered superhard coatings

within the system B-C-N-Ti [1990Kno], [1990Mit]. Improvement of the mechanical properties by

dispersion of B4C particles in a fine grained matrix of TiB2 was reported by [1990Kan]. Creep behavior of

in-situ dual-scale particles-TiB whisker and TiC particulate-reinforced titanium composites was

investigated by [2002Ma]. For superhard materials based on nanostructured high-melting point compounds

see [2001And].

Good thermoelectric properties (354 V/K at 827°C for a specimen with 6 mol% TiB2) were reported for

alloys from the pseudobinary TiB2-B4C system [1998Got].

Miscellaneous

Precipitation of acicular TiB2 from TiC1-x-B alloys containing 0.1 to 1.7 mass% B was observed to increase

microhardness, wear resistance and compressive strength [1979Che]. [1982Evt] studied the interaction of

‘B4C’ with Ti under 0.1 Pa Ar. No interface zones were found in TEM analyses of TiC/TiB2 single layer

and multi-layer coatings, deposited by non-reactive magnetron sputtering on cemented carbide tools

[1991Hil]. The authors of [1980Sha, 1982McC, 1983Shc, 1985Cor, 1991Lev] investigated the process

parameters for the preparation of B-C-Ti alloys by exothermic high temperature Self - Propagating

High-Temperature Synthesis. Thermodynamic analyses of the SHS processes are due to [1999Gor,

2002Mam]. Based on simple thermodynamic calculations of synthesis reactions a processing method

combining conventional melting and combustion synthesis was used to produce Ti-TiB-TiC composites

[1998Ran]. Simple thermodynamic calculations were also applied to boron carbide titanium cermet

synthesis [1986Hal]. [1980Vla] reported on the structure of paramagnetic centres and the formation of

defects in the B-C-Ti system.

References

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73


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B–C–Ti

[1960Por] Portnoi, K.I., Samsonov, G.V., Frolova, K.I., “Some Properties of Alloys of the Titanium

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Titanium Carbide - Boron System Alloys”, Inorg. Mater., 15, 479-482 (1979), translated

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74


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B–C–Ti

[1980Stu] Stubican, V.S., Bradt, R.C., “Directional Solidification of Nonoxide Eutectics”, Report,

ARO-14069.4-MS, Avail. NTIS, from Gov. Rep. Announce Index (U.S.) 1981, 81, 61


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B-C-Ti System”, J. Mater. Sci., 15, 1041-1048 (1980) (Experimental, 9)

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of Cermets in the Titanium - Boron Carbide System” (in Russian), Adgez. Rasplavov i Paika

Mater., 9, 37-39 (1982) (Equi. Diagram, Experimental, 8)

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Self-Sintering of Materials in the System Ti-B-C”, Ceram. Eng. Sci. Proc., 3, 538-554


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Boridy i Silitsidy, Kiev, 97-100 (1982) (Equi. Diagram, Experimental)

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Russian), Fiz. Goreniya Vzryva, 19, 108-111 (1983)

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Boron Carbide - Titanium Boride (B4C-TiB2) System” (in Russian), Sverkhtverd. Mater., 5,


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of B4C Particles”, J. Mater. Sci., 25, 580-584 (1990) (Experimental, 21)

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within the System Ti-B-C-N”, Vacuum, 41, 2184-2186 (1990) (Experimental, 15) see also

Surf. Coat. Technol., 43/44, 107-115 (1990) (Experimental, 12)

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the System Ti-B-C-N”, Surf. Coat. Technol., 41, 351-363 (1990) (Experimental)

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Calculation of Metastable Phase Formation during PVD of Carbide, Nitride, and Boride

Coating”, High Temp. Sci., 27, 295-309 (1990) (Equi. Diagram, Theory, #, 14)

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(Theory, 24)

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[1991Bar] Bartsch, K., Leonhardt, A., Wolf, E., Schönherr, M., “Preparation, Composition and Some

Properties of Codeposited TiB2-TiCx-Coatings”, J. Mater. Sci., 26, 4318-4322 (1991)

(Experimental, Equi. Diagram, Thermodyn., 14)

[1991Gus] Gusev, A.I., “Phase Diagrams for Ordering Systems in the Order-Parameter Functional

Method”, Sov. Phys. Solid State, 32(9), 1595-1599 (1991) (Theory, Equi. Diagram,

Thermodyn., 18), see also Gusev, A.I., “Physical Chemistry of Nonstoichiometric

Refractory Compounds” (in Russian), Chapter 3, Nauka, Moscow, (1991) (Review,

Thermodyn., Crys. Stucture, Equi. Diagram, 102)

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Coatings”, Mat. Sci. Eng., A139, 268-275 (1991) (Experimental, 11)

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609-615 (1991) (Experimental, Equi. Diagram, 19)

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“Structure Formation in Self-Propagating High-Temperature Synthesis of Titanium

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Metall., (3) 82-86 (1991) (Experimental, 12)

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1604-1609 (1991) (Experimental, Crys. Strucrture, 9)

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16(1), 46-60 (1995) (Thermodyn., Equi. Diagram, Review, 36)

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September 1997, Materials Research Center, Royal Institute of Technology, Stockholm,

Sweden., (Thermodyn., Review, 34)

[1997Gus] Gusev, A.I., “Phase Equilibria in the Ternary System Titanium - Boron - Carbon: The

Section TiCy-TiB2 and B4C-TiB2”, J. Solid State Chem., 133(1), 205-210 (1997)

(Thermodyn., Equi. Diagram, 25)

[1998Bit] Bittermann, H., Duschanek, H., Rogl, P., “The Ti-B-C System”, in “Phase Diagrams of

Ternary Metal-Boron-Carbon Systems”, G. Effenberg, (Ed.), ASM-International, MSI,

278-287 (1998) (Review, Crys. Structure, Experimental, Equi. Diagram, 46)

[1998Got] Goto, T., Li, J., Hirai, T., “Thermoelectric Properties of Boron-Rich Boride Composites

Prepared through Eutectic and Peritectic Reactions”, 17th Intl. Conference on

Thermoelectrics, (1998), 574-577 (Experimental, Equi. Diagram, 13)

76


MSIT®

B–C–Ti

[1998Ran] Ranganath, S., Vijayakumar, M., Subrahmanyam, J., “Combustion-Assisted Synthesis of

Ti-TiB-TiC Composite via the Casting Route”, Mat. Sci. Eng., A149, 253-257 (1992)

(Experimental, Thermodyn., 18)

[1999Gor] Gordienko, S. P., “Thermodynamic Analysis of the Reaction of Titanium with Boron

Carbide in a Self-Propagating High-Temperature Synthesis Regime”, Powder Metall. Met.

Ceram., 38, 172-175 (1999) (Experimental, Thermodyn., 3)

[2001And] Andrievski, R.A., “Superhard Materials Based on Nanostructured High-melting Point

Compounds”, Int. J. Refr. Met. Hard Mater., 19, 447-452 (2001) (Mechan. Prop.,

Review, 59)

[2002Ma] Ma, Z.Y., Tjong, S.C., Meng, X.M., “Creep Behaviour of in-situ Dual-scale Particles-TiB

Whisker and TiC Particulate-reinforced Titanium Composites”, J. Mater. Res., 17,

2307-2313 (2002) (Experimental, Mechan. Prop., 26)

[2002Mam] Mamyan, S.S., “Thermodynamic Analysis of SHS Processes”, Key Eng. Mater., 217, 1-8

(2002) (Experimental, Equi. Diagram, Thermodyn., 16)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 [Mas2]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

a = 295.2

c = 469.2

at 25°C [Mas2]

from alloy Ti75B10C15, quenched from

1705°C,

contained also TiB and TiC1-x [1966Rud]

( B)

< 2092

hR333

R3m

B

a = 1093.30

c = 2382.52

a = 1092.2

c = 2381.1

a = 1092.70

c = 2388.65

[1993Wer]

at 1.1 at.% C [1993Wer] linear da/dx, dc/dx

at TiB20 [V-C2]

(C)

< 3827 (S.P.)

hP4

P63/mmc

C-graphite

a = 246.12

c = 670.90

a = 246.023

c = 671.163

a = 246.75

c = 669.78

a = 246.4

c = 672.0

a = 246.4

c = 671.4

[Mas2]

[1967Low]

at 2.35 at.% Cmax (2350°C) [1967Low]

linear da/dx, dc/dx,


2838°C, contained also TiB2 and TiC1-x

[1966Rud]


2636°C, contained also TiB2 [1966Rud]

77


MSIT®

B–C–Ti

‘B4C’

< 2450

hR45

R3m

B13C2

a = 565.1

c = 1219.6

a = 560.7

c = 1209.5

a = 565.1

c = 1219.6

a = 560.7

c = 1209.5

a = 559.8

c = 1212.0

a = 559.7

c = 1212.0

9 to 20 at.% C [1990Ase]

at 91.2 at.% B [V-C2]

at 70.0 at.% B [V-C2]

quenched from 2450°C [1986Ord]

from alloy TiB2+93 mol% B4C,


B25C tP52

P42m

B50C2

a = 872.2

c = 508.0

[V-C2]

also B51C1, B49C3, all metastable?

TiB

< 2190

oP8

Pnma

FeB

a = 610.5

b = 304.8

c = 455.1

a = 612

b = 307.2

c = 456

a = 611

b = 307

c = 456

a = 611.42

b = 305.08

c = 455.90

[V-C2]


1580°C, contained also ( Ti) and TiC1-x

[1966Rud]


1600°C, contained also ( Ti) and TiC1-x

[1966Rud]

from alloy Ti49B46C5, annealed at 1550°C,

contained also TiC1-x and Ti3B4 [1995Dus]

Ti3B4

< 2200

oI14

Immm

Ta3B4

a = 325.9

b = 1373

c = 304.2

a = 326.31

b = 1373.36

c = 303.56

a = 326.30

b = 1372.20

c = 303.84

[V-C2]


contained also TiC1-x and TiB [1995Dus]


contained also TiC1-x and TiB2 [1995Dus]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

78


MSIT®

B–C–Ti

TiB2 hP3

P6/mmm

AlB2

a = 303.1

c = 322.9

a = 303.04

c = 322.94

a = 302.2

c = 322.3

a = 302.4

c = 322.3

a = 302.2

c = 322.4

a = 302.4

c = 322.3

a = 302.8

c = 322.5

a = 302.6

c = 321.3

a = 302.7

c = 321.4

a = 302.

c = 321.3

[V-C2]


contained also Ti2C and Ti3B4 [1995Dus]


2620°C, contained also TiC1-x [1966Rud]


3002°C, contained also TiC1-x [1966Rud]

from alloy Ti27B53C20 quenched from

2712°C, contained also C [1966Rud]

from alloy Ti20B63C17 quenched from

2482°C, contained also C, B4C [1966Rud]

from alloy Ti24B70C6 quenched from 2845°C,

contained also B4C [1966Rud]


from alloy TiB2+3.6 mol% B4C, quenched

from 2759°C [1986Ord]

from alloy TiC+95 mol% TiB2, quenched

from 2610°C [1975Ord]

TiB25 tP52

P42/nnm

TiB25

a = 883.0

c = 507.2

[V-C2]

metastable ?

[1975Amb]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

79


MSIT®

B–C–Ti

TiC1-x cF8

Fm3m

NaCl

a = 432.92

a = 432.60

a = 430.6

a = 432.7

a = 433.2

a = 433.0

a = 432.3

a = 431.12

a = 430.8

a = 432.1

a = 431.0

a = 431.8

a = 432.9

a = 430.9

a = 431.9

a = 432.6

a = 433.0

a = 432.2

a = 432.9

a = 433.0

a = 433.2

a = 432.25

a = 433.0

a = 432.8

a = 433.3

a = 432.2

a = 432.4

TiC0.95, 299 K [V-C2]

TiC0.95, 83 K [V-C2]

TiC0.51 [V-C2]

TiC0.96 [V-C2]

TiC0.95 [1975Ord]

TiC0.8 [1975Ord]

TiC0.68 [1975Ord]


contained TiB and Ti3B4 [1995Dus]


2460°C, contained ( Ti) and TiB2 [1966Rud]


2642°C, contained TiB2 [1966Rud]


2518°C, contained TiB, TiB2 [1966Rud]






2373°C, contained traces ( Ti) [1966Rud]


2800°C, [1966Rud]


2800°C, [1966Rud]


2992°C, [1966Rud]


2517°C, contained TiB2 and B4C [1966Rud]


2625°C, contained TiB2 and B4C [1966Rud]

from alloy TiC0.95+2 mol% TiB2, quenched

from 2930°C [1975Ord]

from alloy TiC0.95+5 mol% TiB2, quenched

from 2900°C [1975Ord]

from alloy TiC0.68+8 mass% TiB2,

linear da/dx [1975Ord]

from alloy TiC0.8+6 mass% TiB2,

linear da/dx [1975Ord]

for TiC0.99B0.01, 1500°C [1975Aiv]

for TiC0.86, 1500°C [1975Aiv]

for TiC0.85B0.1, 1500°C [1975Aiv]

for TiC0.81B0.17, 1500°C [1975Aiv]

Ti1.86B48C2 tP52

P42/nnm

TiB25

a = 887.6

c = 506.2

[V-C2] [1980Amb]

metastable (?)

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

80


MSIT®

B–C–TiT

ab

le 2

: I

soth

erm

al R

eact

ion

s in

the

B-C

-Ti

Sy

stem

; C

om

par

ison

of

Cal

cula

ted D

ata

wit

h E

xper

imen

tal

Dat

a fr

om

[1

966

Ru

d]

Cal

cula

ted

Dat

aE

xp

erim

enta

l D

ata

T [

K (

°C)]

Rea

ctio

nT

yp

eT

[K

(°C

)]R

eact

ion

Ty

pe

29

14

(26

41)

LT

iB2

+T

iC1

-x

e 2(m

ax)

28

93 (

26

20)

LT

iB2

+T

iC1

-x

e(m

ax)

at.%

Ti

43.9

6

33

.33

53

.96

at.%

Ti

45

~

34

5

3

at.%

B3

3.9

9

66

.67

3.1

9

at

.% B

31

~

63

~

6

at.%

C2

2.0

5

0.0

0

42

.85

at.%

C2

4

<3

4

1

27

36

(24

63)

LT

iB2

+C

e 3

(max

)27

80 (

25

07)

LT

iB2

+C

e(

max

)

at.%

Ti

22.2

8

33

.33

0.0

0

at

.% T

i2

3

~3

4

<1

at.%

B4

4.7

7

66

.67

0.6

9

at

.% B

45

~

64

~

2

at.%

C3

2.9

5

0.0

0

99

.31

at.%

C3

2

<2

>

97

26

73

(24

00)

LT

iB2

+T

iC1

-x

+ C

E1

26

73 (

24

00)

LT

iB2

+T

iC1

-x+

CE

at.%

Ti

29.1

7

33

.33

51

.08

0.0

0

at.%

Ti

29

3

4

52

<1

at.%

B3

3.4

9

66

.67

1.5

20.2

8

at.%

B3

7

64

2

<2

at.%

C3

7.3

4

0.0

0

47

.40

99

.72

at

.% C

34

<

24

6

>97

26

39

(23

66)

LT

iB2

+‘B

4C

’

e 5 (

max

)25

83 (

23

10)

LT

iB2

+‘B

4C

’

e (m

ax)

at.%

Ti

7.9

4

33

.33

0.0

0

at

.% T

i~

5

>3

3

<1

at.%

B7

8.5

2

66

.67

82

.23

at.%

B8

0

~6

5

~8

2

at.%

C1

3.5

4

0.0

0

17

.77

at.%

C~

15

<

2

~1

7

25

40

(22

67)

LT

iB2

+‘B

4C

’ +

CE

225

13 (

22

40)

LT

iB2

+‘B

4C

’+

CE

at.%

Ti

10.0

1

33

.33

0.0

00.0

0

at.%

Ti

~1

0

~3

4

<1

<1

at.%

B6

3.9

2

66

.67

80

.40

1.9

4

at.%

B~

62

~

64

~

80

~3

at.%

C2

6.0

7

0.0

0

19

.60

98

.06

at

.% C

~2

8

<2

>

19

>96

23

90

(21

17)

L+

TiB

2T

i 3B

4 +

TiC

1-x

U1*

..

.

...

..

....

...*

at.%

Ti

64.2

8

33

.33

42

.86

62

.18

at

.% T

i..

.

...

..

....

at.%

B3

1.1

5

66

.67

57

.14

3.2

9

at.%

B..

.

...

..

....

at.%

C4

.57

0.0

0

0.0

034

.53

at

.% C

...

..

.

...

...

81


MSIT®

B–C–TiT

ab

le 2

: (c

on

tinued

)

*)

See

sec

tion

“In

var

iant

Eq

uil

ibri

a”

Cal

cula

ted

Dat

aE

xper

imen

tal

Dat

a

T [

K (

°C)]

Rea

ctio

nT

ype

T [

K (

°C)]

Rea

ctio

nT

yp

e

23

82

(21

09)

L+

Ti 3

B4

TiB

+T

iC1

-xU

2*

24

33

(2

160

)L

+T

iB2

TiB

TiC

1-x

U*

at.%

Ti

64.5

0

42

.86

50

.00

62

.27

at

.% T

i~

59

~

34

~

50

~6

4

at.%

B3

1.0

0

57

.14

50

.00

3.2

7

at.%

B~

34

>

64

>

48

~4

at.%

C4

.50

0.0

0

0.0

034

.46

at

.% C

~7

<

2

<2

~3

2

23

58

(20

85)

L+

‘B4C

’ T

iB2

+ (

B)

U3*

22

89

(2

016

)L

+T

iB2

‘B4C

’(

B)

E*

at.%

Ti

2.1

6

0.0

0

33

.33

0.0

0

at.%

Ti

~3

3

<1

<1

at.%

B9

7.3

3

90

.15

66

.67

98

.58

at

.% B

~97

~

66

>

88

~9

9

at.%

C0

.51

9.8

5

0.0

01.4

2

at.%

C~

1.5

~

1

~11

<1

18

07

(15

34)

L(

Ti)

+T

iB +

TiC

1-x

E3

17

83

(1

510

)L

(T

i)+

TiB

TiC

1-x

E

at.%

Ti

91.3

4

98

.03

50

.00

66

.95

at

.% T

i93

>

98

~

51

~6

8

at.%

B6

.95

0.2

0

50

.00

1.2

2

at.%

B5

~

1

>4

8~

2

at.%

C1

.71

1.7

7

0.0

031

.83

at

.% C

2

~1

<

1~

30

11

93

(92

0)

(T

i)+

TiC

1-x

+T

iB

(T

i) P

1*

11

63

(8

90)

(T

i)+

TiC

1-x

TiB

+ (

Ti)

U*

at.%

Ti

99.5

6

63

.78

50

.00

98

.29

at

.% T

i...

..

.

...

...

at.%

B0

.13

0.2

5

50

.00

0.1

5

at.%

B...

..

.

...

...

at.%

C0

.31

35

.97

0.0

01.5

6

at.%

C...

..

.

...

...

82


MSIT®

B–C–Ti

Fig

. 1

:

B-C

-Ti.

Rea

ctio

n s

chem

e

l T

iC1-x

+ C

27

72

e 1

lβT

i +

TiC

1-x

16

54

e 7

βTi

+ T

iC1-x

αTi

91

9p

4

l 'B

4C

' + C

23

90

e 4

l +

'B4C

'(β

B)

21

03

p3

L T

iB2

+ T

iC1-x

+ C

24

00

E1

L T

iB2

+ T

iC1

-x

26

41

e 2(m

ax)

L T

iB2

+ C

24

63

e 3(m

ax)

L T

iB2

+ 'B

4C

'

23

66

e 5(m

ax)

L T

iB2

+ 'B

4C

' + C

22

67

E2

L +

TiB

2 T

i 3B

4+

TiC

1-x

21

17

U1

L +

Ti 3

B4

TiB

+ T

iC1

-x2

109

U2

L +

'B4C

' T

iB2

+(β

B)

20

85

U3

L(β

Ti)

+ T

iB +

TiC

1-x

15

34

E3

(βT

i)+

TiB

+T

iC1-x

(αT

i)9

20

P1

l +

TiB

2 T

i 3B

4

21

99

p1

l +

Ti 3

B4

TiB

21

84

p2

l T

iB2

+(β

B)

20

60

e 6

l(β

Ti)

+ T

iB

15

41

e 8

(βT

i) +

TiB

(α

Ti)

88

3p

5

(βT

i)+

TiB

+T

iC1

-x

TiB

2+

'B4C

'+(β

B)

L+

TiB

2+

(βB

)

L+

TiB

+T

iC1

-x

TiB

2+

'B4C

'+C

L+

Ti 3

B4+

TiC

1-x

TiB

2+

Ti 3

B4+

TiC

1-x

Ti 3

B4+

TiB

+T

iC1

-x

(αT

i)+

TiC

1-x

+T

iB

TiB

2+

TiC

1-x

+C

C-T

iB

-CB

-C-T

iB

-Ti

83


MSIT®

B–C–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti C

B Data / Grid: at.%

Axes: at.%

e1, 2772

e6, 2060

e7,1654

e4, 2390

p3, 2103

'B4C'

2452

TiC1-x

3084

TiB2

3216

4000

3800

3600

3200

28002600

30002800

1600

1600

1800

2000

p2, 2184

2600

3000

2800

26002400

2200

maximum in liquid trough

four-phase equilibrium

e5

(max), 2366

E2, 2267

e3(max), 2463

E1, 2400

e2(max), 2641

U1, 2117

U2, 2109

U3, 2085

p1, 2199

e8, 1541

E3, 1534

TiB

Ti3B

4

Experimental data from [1966Rud]:

Fig. 3: B-C-Ti.

Calculated liquidus

projection and

isothermal reactions

20

40

60

80

20 40 60 80

20

40

60

80

Ti C

B Data / Grid: at.%

Axes: at.%

e,2776

E,2016

e,2080(βB)

e,2080

e,2310

E,2240

e,2507

E,2400

e, 2620

U,2160

E,1510

e,1650

e,1540

p,2180

TiB

TiC1-x

e,2380

(βTi)

C

TiB2

'B4C'

Fig. 2: B-C-Ti.

Projection of liquidus

surface; from

[1966Rud]

84


MSIT®

B–C–Ti

20 40 60 80

2000

2250

2500

2750

3000

3250

Ti 33.33

B 66.67

C 0.00

C

C, at.%

Te

mp

era

ture

, °C

L+C

L+TiB2

TiB2+C

2507±15°C

3225±20°C

32±2<2

10 20 30 40

2250

2500

2750

3000

3250

Ti 33.00

C 0.00

B 67.00

Ti 52.00

C 48.00

B 0.00C, at.%

Te

mp

era

ture

, °C

L

L+TiB2

L+TiC1-x

TiB2+TiC1-x

<1 27.4±1 ~45.6

3225±20°C

2620±15°C

Fig. 4: B-C-Ti.

Experimentally

derived isopleth

TiB2 - C; after

[1966Rud]

Fig. 5: B-C-Ti.

Experimentally

derived isopleth

TiB2 - TiC1-x; from

[1966Rud]

85


MSIT®

B–C–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti C

B Data / Grid: at.%

Axes: at.%

TiB2

Ti3B

4

TiB

TiC1-x

'B4C'

(βTi)+TiB+TiC1-x

TiB2+TiC

1-x+C

TiB2+C+'B

4C'

TiB

2+B

+'B

4C

'

(βTi)

(βB)

TiB+Ti3B

4+TiC

1-x

Ti3B

4+TiC

1-x+TiB

2

10

2000

2250

2500

2750

3000

3250

Ti 33.00

B 67.00

C 0.00

Ti 0.00

B 81.70

C 18.30C, at.%

Te

mp

era

ture

, °C

<0.4 16

L

~2450°C

3225°C

2310

TiB2+L

TiB2+B4C

Fig. 7: B-C-Ti.


section at 1400°C

Fig. 6: B-C-Ti.

Experimentally

derived isopleth

TiB2 - 'B4C'; after

[1966Rud]

86


MSIT®

B–C–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti C

B Data / Grid: at.%

Axes: at.%

TiB2

Ti3B

4

TiB

TiC1-x

'B4C'

TiB2+TiC

1-x+C

TiB2+C+'B

4C'

L+T

iB2+'B

4C

'

L

L

L+T

iB2+T

iC1-x

L+TiB+Ti3B

4

20

40

60

80

20 40 60 80

20

40

60

80

Ti C

B Data / Grid: at.%

Axes: at.%

TiB2

TiC1-x

'B4C'

TiB2+TiC

1-x+C

TiB2+L+'B

4C'

TiB 2+L+'B 4

C'

L

L

L+T

iB2+T

iC1-x

TiB2+L+C

L+'B

4 C'+

C

L

Fig. 8: B-C-Ti.


section at 2150°C

Fig. 9: B-C-Ti.


section at 2300°C

87


MSIT®

B–C–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti C

B Data / Grid: at.%

Axes: at.%

TiB2

TiC1-x

'B4C'

TiB2+TiC

1-x+C

L

L+

TiB

2+

TiC

1-x

TiB2+L+C

L

20 40 60 80

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

Ti 50.00

B 50.00

C 0.00

C

C, at.%

Te

mp

era

ture

, °C

L+C

L

L+TiC1-xL+TiB2

TiB2+TiC1-x

TiB2+TiC1-x+C

TiB+Ti3 B4+TiC1-x

Ti3 B4+TiB2+TiC1-x

2400

2109°C2117

Fig. 10: B-C-Ti.


section at 2400°C

Fig. 11: B-C-Ti.

Calculated isopleth

Ti0.5B0.5 - C

88


MSIT®

B–C–Ti

20 40 60 80

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

Ti 50.00

B 0.00

C 50.00

B

B, at.%

Te

mp

era

ture

, °C

L

L+TiB2L+TiC1-x

TiB2+TiC1-x+C

TiB2+C

TiB2+C+'B4C'

TiB2+'B4C'

TiB2+(βB)+'B4C'

L+'B4C'

L+(βB)TiC1-x+C

2400 2267

2085

(βB)

20 40 60 80

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

Ti Ti 0.00

B 81.70

C 18.30B, at.%

Te

mp

era

ture

, °C

L

L+TiB2

L+TiB2+TiC1-x

L+TiB2+C L+'B4C'

L+TiB2+'B4C'

TiB2+'B4C'

TiB2+C

TiC1-x+TiB2+C

TiC1-x+TiB2

L+TiC1-x

2117

1534

920

2267

2400

TiC

1-x+

Ti

3B4+

TiB

2

TiB

+T

iC1-x+

Ti

3B4

TiB+TiC1-x

TiC1-x+Ti3B4

(αTi)+TiB (αTi)+TiB+TiC1-x

(αTi)+(βTi)+TiB

(βTi)+TiB+TiC1-x

(βTi)+TiB

L+TiB(βTi)+L

(αTi)

(βTi)

2109

TiB2+C+'B4C'

Fig. 12: B-C-Ti.

Calculated isopleth

Ti0.5C0.5 - B

Fig. 13: B-C-Ti.

Calculated isopleth

from Ti to ’B4C’

89


MSIT®

B–C–Ti

20 40 60 80

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

Ti 56.16

B 0.00

C 43.84

Ti 0.00

B 81.70

C 18.30B, at.%

Te

mp

era

ture

, °C

L

L+'B4C'

TiC1-x+L

24002267TiB2+C+L

TiB2+L

TiB2+'B4C'+L

TiB2+'B4C'

TiB2+C

TiC1-x+TiB2

TiC1-x+TiB2+C TiB2+C+'B4C'

TiC1-x+TiB2+L

TiC1-x

Fig. 14: B-C-Ti.

Calculated isopleth

from ’TiC1-x’ to

’B4C’

90


MSIT®

B–Li–N

Boron – Lithium – Nitrogen

Oksana Bodak

Literature Data

The literature data up to 1987 were compiled and critically reviewed by [1992Rog, 1995Pav]. The only

ternary compound Li3BN2 forms in the B-Li-N system. Li3BN2 appears in four polymorphic modification,

two of which are high-pressure modifications. The ternary compound Li3BN2 was first identified by

[1961Gou] who gave its melting point as 870 5°C and measured its density as 1.755 g cm-3. Chemical

analysis of the product obtained by reacting Li3N and BN in a nitrogen atmosphere at 650 to 750°C for 2 h

gave Li3.05BN1.87 (after normalizing the 94.86 mass% analysis to 100%). This compound was designated

as Li3BN2 (low temperature modification) by [1987Yam] who prepared it by heating compressed Li3N

and BN powder mixtures in the molar ratios of 1.0 to 1.2 at 1027°C for 20 min, cooling at 2°C min-1 to

727°C and subsequently more rapidly to room temperature. [1987Yam] also studied the transformation of

Li3BN2 to Li3BN2 by DTA and X-ray diffraction analysis of quenched samples held at temperatures

from 700 to 900°C. The polymorphic transformation is observed at about 887°C by DTA and at about

862°C from examination of quenched samples; Li3BN2 melts at 916°C [1987Yam]. The melting

temperature of 870 5°C reported by [1961Gou] for Li3BN2 is likely to correspond to Li3BN2 Li3BN2

phase transformation [1992Rog]. [1986Yam] investigated the crystal structure of Li3BN2 single crystals

(high temperature modification) that were chemically analyzed as Li3BN1.96 (normalized from 98.5 mass%

analysis to 100 mass%), while the structural charaterization was obtained by X-ray diffraction analysis.

Experiments were done using the mixtures Li3N/BN in molar ratio 1.0 to 1.2, which were heated above the

melting point of the product at 927°C, kept at this temperature for 7h and cooled to 727°C at a rate of

1.5-3°C/h under a nitrogen stream.

[1969DeV, 1972DeV] investigated the high-temperature - high-pressure phase transformations of Li3BN2

in the pressure range of 1 to 6.5 GPa (10 to 65 kbar) and the temperature range of 300 to 1900°C. [1969DeV]

carried out experiments in a pyrophyllite-Ta cell in a belt apparatus. Pressure was applied to the cell and

then the temperature was raised at a rate of about 400°C min-1; both p and T were held for 10 to 15 min, the

samples were quenched in about 30 sec to the room temperature with still applied pressure. Li3BN2 (I) was

recognized to be identical with phase lebelled as “complex” by [1961Wen]. Phases were identified by room

temperature X-ray diffraction analysis and by transmitted light optical microscopy. The temperature of

1550°C corresponds to the lowest temperature for the growth of cubic BN from metallic Li [1966Kud].

Binary Systems

The binary systems are accepted from [Mas2], except the B-N system that is accepted from [2003Rec]. Data

of [1984Mai] on the crystal structure of the new lithium borides are includeds in Table 1.

Solid Phases

Ternary solid phases identified in the B-Li-N system are shown in Table 1, as well as some solid phases

from boundary binary systems. p-T phase diagram for Li3BN2 is given in Fig. 1. The Li3BN2 (I)- Liquid

equilibrium line is probably metastable because [2002Tur] observed an incongruent melting of Li3BN2 (I)

around 1350°C under 5.3 GPa. Li3BN2 is a low-pressure phase [1986Yam]. The high-pressure phase,

Li3BN2 (I), is well known [1969DeV, 1972DeV]. A second high-pressure phase, designated Li3BN2 (II)

must exist, given the change in the curvature of the Li3BN2 (I) phase region at temperatures below 800°C,

Fig. 1. [1969DeV] found no difference between the diffraction patterns of Li3BN2 and Li3BN2 (I); they

suggest the need to use high pressure X-ray diffraction techniques to clarify this portion of p-T diagram. The

X-ray diffraction patterns given by [1961Gou] and [1969DeV] agree with that published by [1987Yam] for

Li3BN2. The X-ray diffraction patterns of a high-temperature - high-pressure phase observed by

[1961Wen] agrees with that given for Li3BN2 (I) by [1969DeV].

91


MSIT®

B–Li–N


Two different diagrams have been proposed for the Li3N-BN section. [1972DeV] studied the Li3BN2

(I)-BN section at a pressure of 5 GPa (50 kbar) and eutectic was observed at 49 at.% Li and 1610 15°C

between Li3BN2 (I) and cubic BN. More recently, [2002Tur] showed that Li3BN2 melted incongruently

around 1350°C under 5.3 GPa and observed that the temperature of the peritectic equlibrium was the lowest

temperature of the cBN synthesis in the B-Li-N system under 5.3GPa.

This apparent contradiction between both diagrams is well explained with the hypotheses of two

equilibrium diagrams. The diagram proposed by [2002Tur] is a stable one whereas that proposed by

[1972DeV] is a metastable. The existence of metastable Li3BN2 under 5.3 GPa at temperature higher that

1350°C agrees with the observed nucleation and growth of cubic BN by using Li3BN2 as a solvent-catalyst

[1992Nak, 1993Boc].

Figure 2 shows both stable and metastable diagrams for the Li3N-BN section under 5.3 GPa. This figure has

been modified to take into account the true transition temperature of cBN hBN given around 2950°C by

[2002Tur]. The accepted transition temperature lies around 1770°C under 5.3 GPa [1966Kud, 1972DeV,

1993Nak].

Thermodynamics

The standard enthalpy of formation of Li3BN2 was determined by drop solution calorimetry [1999McH].

Li3BN2 was synthesized from a mechanical mixture of Li3N and BN (Aldrich, 99.9%) following the

procedure of [1961Gou]. The mixture of binaries was heated at 800°C for 24 h. The gray product was

identified as nearly single phase -Li3BN2 via comparison of its powder XRD pattern to that reported by

[1987Yam]. Nitrogen analysis of the product yielded 45.96 0.46% N (theoretical = 46.96%), which

corresponds to a 2.1 1.0 mass% LiOH impurity. High-temperature oxide melt drop solution calorimetry

was performed in a Tian-Calvet twin microcalorimeter.

The standard enthalpy of formation of Li3BN2 [1999McH, 2001Nav] measured from the elements is

fH° = -534.5 16.7 kJ mol-1. Using previously determined fH°(Li3N) = -166.1 4.8 kJ mol-1 and

fH°(BN) = -250.91 kJ mol-1 listed in JANAF tables, [1999McH] calculated rH° = -117.5 17.5 kJ mol-1

for the reaction Li3N+BN Li3BN2 at 298 K.

Miscellaneous

Li3BN2 has been proposed as a catalyst for the production of cubic BN (known as “borazon”) from

hexagonal BN [1992Nak, 1993Nak]. Li3BN2 plays the role of a solvent of BN at high temperature

(>1550°C) and high pressure (>5 GPa). [1993Boc] observed that cubic BN crystals formed in the Li3N-BN

system were generally more perfect than with Ca3N2 or Mg3N2. The phase boundary between hexagonal

BN and cubic BN was experimentally determined at p = T/200-4.9 (p in GPa; T in K) [1992Nak].

References

[1961Gou] Goubeau, J., Anselment, W., “Ternary Metal Boronitrides.” (in German), Z. Anorg. Allg.

Chem., 310, 248-260 (1961) (Equi. Diagram, Experimental, 14)

[1961Wen] Wentdorf, R.H., Jr., “Synthesis of the Cubic Form of Boron Nitride”, J. Chem. Phys., 34(3),


[1966Kud] Kudaka, K., Konno, H., Matoba, T., “Parameters for the Crystal Growth of Cubic Boron

Nitride” (in Japanese), Kogyo Kogaku Zasshi, 69, 365-369 (1966) (Crys. Structure,

Thermodyn., Experimental, 14)

[1969DeV] DeVries, R.C., Fleischer, J.F., “The System Li3BN2 at High Pressures and Temperatures”,

Mater. Res. Bull., 4, 433-441 (1969) (Equi. Diagram, Crys. Structure, Experimental, 9)

[1972DeV] DeVries, R.C., Fleischer, J.F., “Phase Equilibria Pertinent to the Growth of Cubic BN”,

J. Cryst. Growth, 13-14, 88-92 (1972) (Equi. Diagram, Experimental, 9)

[1984Mai] Mair, G., “On the Lithium-Boron System” (in German), Ph.D. Thesis, University of

Stuttgart, pp. 97 (1984) (Equi. Diagram, Crys. Structure, Experimental, 57)

92


MSIT®

B–Li–N

[1986Yam] Yamane, H., Kikkawa, S., Oriuchi, H., Koizumi, M., “Structure of a New Polymorth of Li-B

Nitride, Li3BN2”, J. Solid State Chem., 65, 6-12 (1986) (Crys. Structure, Experimental, 25)

[1987Yam] Yamane, H., Kikkawa, S., Koizumi, M., “High- and Low-Temperature Phase of Li-B

Nitride, Li3BN2: Preparation Phase Relation, Crystal Structure, and Ionic Conductivity”,

J. Solid State Chem., 71, 1-11 (1987) (Crys. Structure, Experimental, 14)

[1992Nak] Nakano, S., Fukunaga, O., “New Concept of the Synthesis of Cubic Boron Nitride”, Recent

Trends in High Pressure Science and Technology, XIII AIRAPT - Intl. Conf. on High

Pressure Technology, 1991, 687-691 (1992) (Equi. Diagram, Experimental, 14)

[1992Rog] Rogl, P., Schuster, J.C., “Phase Diagrams of Ternary BN and SiN System” Monogr.Ser. of

Alloy Phase Diag., 52-55 (1992) (Equi. Diagram, Crys. Structure, Review, 13)

[1993Boc] Bocquillon, G., Loriers-Susse, C., Loriers, J., “Synthesis of Cubic Boron Nitride Using Mg

and Pure or M’-Doped Li3N, Ca3N2 and Mg3N2 with M’=Al, B, Si, Ti”, J. Mater. Sci., 28,


[1993Nak] Nakano, S., Ikawa, H., Fukunaga, O., “Synthesis of Cubic Boron Nitride Using Li3BN2,

Sr3B2N4 and Ca3B2N4 as Solvent-Catalysts”, Diamond and Related Materias, 3, 75-83



of Carbon-Doped -Rhombohedral Boron”, Phys. Status Solidi (B), B179, 489-511 (1993)


[1995Pav] Pavlyuk, V., Bodak, O., “Boron-Lithium-Nitrogen”, MSIT Ternary Evaluation Program, in



Assessment, 7)

[1999McH] McHale, J.M., Navrotsky, A., DiSalvo, F.J., “Energetics of Ternary Nitride Formation in the

(Li,Ca)-(B,Al)-N System”, Chem. Mater., 11, 1148-1152 (1999) (Thermodyn.,

Experimental, 29)

[2001Nav] Navrotsky, A., “Thermochemical Studies of Nitrides and Oxynitrides Oxidative Oxide Melt

Calorimetry”, J. Alloys Compd., 321, 300-306 (2001) (Thermodyn., Review, 28)

[2002Tur] Turcevic, V.Z., “Phase Diagram and Synthesis of Cubic Boron Nitride”, J. Phys.: Condens.

Matter, 14(44), 10963-10968 (2002) (Equi. Diagram, Experimental, Calculation, 16)

[2003Rec] Record, M.Ch., Tedenac, J.-C., “B-N (Boron-Nitrogen)”, MSIT Evaluation Program, in


GmbH, Stuttgart, to be published, (2002) (Equi. Diagram, Review, 50)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Li)

< 180.6

cI2

Im3m

W

a = 351.0 pure Li at 25°C

[V-C2]

( B)

< 2092

hR333

R3m

B

a = 1093.30

c = 2382.52

[Mas2, 1993Wer]

( N)

< -237.54

cP8

Pa3

N

a = 566.1 [2003Rec]

LiB12 tP216

P41212

AlB12

a = 1016

c = 1428

[V-C2, Mas2]

93


MSIT®

B–Li–N

LiB6 - - [Mas2]

Li3B14 tI136

I42d

Li3B14

a = 1076.4 0.5

c = 894.7

[1984Mai, V-C2, Mas2]

LiB4 c** a = 720 [Mas2]

Li6B19 hP51

P6/mmm

Li6B19

a = 823.3

c = 415.6

[1984Mai]

LiB3 tP16

P4/mbm

Li2B6

a = 597.8

c = 418.8

[1984Mai]

LiB2 - - [Mas2]

LiB t** a = 1391

c = 815

[Mas2]

Li5B4

(Li7B6)

hR27

R3m

Li5B4

a = 697

c = 854

= 90°

[V-C2, Mas2]

Li3B c** - [Mas2]

LiN3 mC8

C2/m

AuSe

a = 562.7

b = 331.9

c = 497.9

= 107.4°

[V-C2, Mas2]

Li3N

< 815

hP4

P6/mmm

Li3N

a = 364.8

c = 387.5

[V-C2, Mas2]

melts at 1017°C under 5.3 MPa

[2002Tur]

wBN hP4

P63/mmc

ZnS

a = 255.0 0.5

c = 423 1

[2003Rec]

cBN cF8

F43m

ZnS

a = 361.53 [2003Rec]

hBN

< 2397

hP4

P63mc

BN

a = 250.4

c = 666.1

[2003Rec]

rBN hR6 a = 250.4

c = 999.1

[2003Rec]

Compressed hBN mC24

C2/c or Cc

a = 433

b = 250

c = 310 to 330

= 92-95°

[2003Rec]

B25N tP62

P42m

B25N

a = 863.4 0.4

c = 512.8 0.3

[2003Rec]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

94


MSIT®

B–Li–N

* Li3BN2

< 862

tP12

P42212

Li3BN2

a = 464.35

c = 525.92

[1987Yam]

m=1.75g cm-3

* Li3BN2

916 - 862

mP24

P21/c

Li3BN2

a = 515.02

b = 708.24

c = 679.08

= 112.956°

[1986Yam, 1987Yam]

=1.74 g cm-3

x=1.737 g cm-3

* Li3BN2 (I) - - [1961Wen, 1969DeV, 1972DeV],

HP-HT

* Li3BN2 (II) - - [1969DeV], HP-HT

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

Temperature, °C

Pre

ssu

re,

GP

a

1200

0

1600800400

2

4

6

LαLi BN3 2

Li BN (II)3 2

Li BN (I)3 2

Fig. 1: B-Li-N.

p-T diagram for

Li3BN2 [1969DeV]

95


MSIT®

B–Li–N

10 20 30 40

750

1000

1250

1500

1750

Li 75.00

B 0.00

N 25.00

Li 0.00

B 50.00

N 50.00B, at.%

Te

mp

era

ture

, °C

L+cBN

Li3BN2+cBN

L+Li3BN2

Li3N+Li3BN2

850°C

1350°C

1770°C

L+hBN

L

Li3BN2

1650°C

Fig. 2: B-Li-N.

The Li3N - BN

quasibinary section

at 5.3GPa.

In dashed lines:

metastable diagram

96


MSIT®

B–Mg–N

Boron – Magnesium – Nitrogen

Peter Rogl, Andreas Leithe-Jasper, Takaho Tanaka

Literature Data

Industrial high-pressure production of super-hard cubic boron nitride (cBN), also called “borazon”, relies

on the use of Mg, Mg3N2 and ternary magnesium boron nitrides as catalysts/solvents. Based on preliminary

high-pressure data [1960Wen, 1966Kud, 1967Fil], defining the minimum conditions for the formation of

cubic BN in the system B-Mg-N at ~5.5 GPa and ~1700°C in the presence of magnesium borides, numerous

high-pressure, high-temperature investigations essentially concern the subsystems Mg-BN [1960Wen,

1966Kud, 1967Fil, 1971Ush, 1972Vri, 1975Fuk, 1979End1, 1979End2, 1991Bin1, 1991Bin2, 1993Boc]

and Mg3N2-BN [1960Wen, 1979End2, 1981Ely, 1981Sat, 1986Yam, 1988Lor1, 1988Lor2, 1989Hoh,

1992Nak1, 1992Nak2, 1993Nak1, 1993Boc, 1994Gla, 1994Zhu, 1995Lor1, 1995Lor2, 1997Lor,

1999Kul2]. Most of the studies are devoted to derive optimum conditions for the catalyst/solvent-dependent

conversion of hBN to cBN and therefore result in a pressure versus temperature (pT)-diagram with stability

regions for hBN and cBN [1966Kud, 1971Ush, 1972Vri, 1975Fuk, 1981Sat, 1981Ely, 1986Yam,

1991Bin2, 1992Nak1, 1992Nak2, 1993Nak1, 1993Boc]. A smaller number of investigations is directly

concerned with the constitution of the subsystems Mg-Mg3BN3 ( 1) (earlier given as “Mg-Mg3B2N4”

[1979End2]), Mg3N2-Mg3BN3 [1992Lor, 1993Lor, 1993Nak1], BN-Mg3BN3 (earlier

“BN-Mg3B2N4( 3)”) [1979End2, 1993Nak1, 1995Lor1, 1995Lor2, 1997Lor] and BN-MgB2 [1999Sol]. In

these cases phase relations are presented in temperature versus composition diagrams derived at a fixed

pressure. The pressure dependency of three (metastable) eutectic reactions L Mg3N2+hBN,

L Mg3BN3+hBN and L “Mg3B2N4”+hBN was claimed from in-situ DTA measurements and X-ray

diffraction on the quenched products [1994Gla].

High pressures were produced in simple cylinder-piston apparatus [1979End2], in belt-type equipment

[1979End1, 1981Sat, 1988Lor2, 1989Hoh, 1993Boc, 1995Sin2], in a multi-anvil press [1991Bin1,

1991Bin2, 1992Nak1, 1993Nak1, 1992Lor, 1993Lor, 1995Lor1, 1995Lor2, 1997Lor, 2003Lee, 2003Kar]

or in a toroidal press [1981Ely, 1994Gla, 1994Zhu, 1999Kul1, 1999Kul2, 2000Kul]. Experiments to all

these investigations were generally conducted in internally heated high pressure cells in which the

experimental mixture was contained in a Ta liner inside a BN sleeve [1972Vri] or (Mg, Mg3N2) and

mixtures of (Mg3N2+hBN) powders in a BN-sleeve inside a graphite heater [1979End1, 1979End2,

1981Sat, 1989Hoh, 1993Boc, 1995Lor2]. In some cases (Mg+hBN) powders were directly charged in a

graphite heater with Mo-end discs [1991Bin1, 1991Bin2] or sample mixtures were enclosed in Mo- or Ta-

foil and heated by a carbon heater through a hBN or pyrophyllite insulation sleeve [1992Nak1, 1999Kul1,

1999Kul2, 2000Kul]. [1994Zhu] used sealed ampoules made of heat resistant steel. [1993Nak1, 1993Nak2]

contained the sample mixtures in Pt-foil within a NaCl pressure medium inside a graphite heater. The high

pressure high temperature runs usually were pursued according to the following procedure: after raising the

pressure to the pre-designed value, the sample is heated slowly to the predetermined temperature points.

After a reaction time of 15 to 30 min, the sample is rapidly cooled to room temperature whereupon the

pressure was released. Temperature measurements (estimated uncertainty 30°C) were performed using

Pt/PtRh thermocouples with generally no corrections being made for pressure. Pressure calibration was

made using the change in resistivity of pure metal standards (estimated accuracy of measurements about

3%). More sophisticated experimental arrangements were used for in-situ synchrotron experiments

[1986Yam, 1992Lor, 1993Lor, 1995Lor1, 1995Lor2, 1997Lor, 1999Sol] as well as for the in-situ DTA

experiments at fixed pressure [1979End2] (at 2.5 GPa, heating/cooling 35 to 40°C/min), [1989Hoh] (at 6.5

GPa), [1994Gla] (3.0 to 8.0 GPa).

The reaction products were examined by light optical microscopy and X-ray diffraction [1960Wen,

1966Kud, 1967Fil, 1971Ush, 1972Vri, 1975Fuk, 1979End1, 1979End2, 1981Sat, 1981Ely, 1986Yam,

1988Lor2, 1989Hoh, 1991Bin1, 1991Bin2, 1992Nak1, 1992Nak2, 1993Nak1, 1993Boc, 1994Gla,

1994Zhu, 1995Lor1, 1995Lor2, 1995Sin2, 1997Lor, 1999Kul1, 1999Kul2, 2000Kul, 2002Mir] and in some

97


MSIT®

B–Mg–N

cases by infrared spectroscopy [1971Ush, 1975Fuk, 1979End2, 1981Ely, 1999Kul1, 1999Kul2, 1999Kul3,

2000Kul]; the crystal size and crystal habit was checked by SEM [1966Kud, 1971Ush, 1975Fuk, 1979End1,

1981Sat, 1988Lor2, 1991Bin1, 1991Bin2, 1992Nak1, 1992Nak2, 1993Nak1, 1993Boc, 1995Sin2].

[1981Ely, 1989Pik] employed chemical analysis; for the thermal stability of the products gas

chromatography and derivatography was used [1981Ely]. To obtain cBN crystals, residual catalyst and hBN

were removed by leaching in diluted HCl and subsequently in NaF-H2SO4 solution [1979End1, 1981Sat,

1992Nak2, 1991Bin1, 1991Bin2, 1993Boc, 1995Lor1, 1995Lor2, 1997Lor], in HF-HNO3 solution

[1993Nak1] or in molten hydroxides [1988Lor2]. Separation was also achieved by flotation in bromoform

[1993Boc].

Starting materials were in all cases Mg-flakes with a minimum purity of >99.5 mass%. Information how to

prepare Mg3N2 is presented by [1981Sat, 1988Lor2, 1993Nak1, 1999Kul1]. Oxygen levels in Mg3N2 were

given as 0.6 mass% [1989Hoh] and 0.8 mass% [1988Lor2]. Hexagonal boron nitride was either used

directly in form of commercial powder or from hot-pressed material. In some cases its graphitization index

(GI) was given. Oxygen levels were determined form neutron activation: [1979End1] (10-15 m, 1.9(1) and

7.9(4) mass% O, GI=1.56 and 1.39, respectively), [1979End2] (hot-pressed and dried under 150°C at 10 Pa;

1.9 mass% O), [1981Sat] (various grades of hot-pressed and powder BN from 0.3 to 8.8 mass% O and 24.1

mass% O), [1989Hoh] (0.6 mass% O, GI=1.6) or from ESCA: [1991Bin1] (powder, 5-12 m, 51 at.% O,

32.7 at.%N and 7.8 at.% C, GI=1.68; from rod, 0-2 m, 37.7 at.% O, 48.7 at.% N and 6.6 at.% C, GI=1.79),

[1991Bin2] (powder from rod, 0-2 m, 37.7 at.% O and 48.7 at.% N, GI=1.79). Without details on the

method of determination, oxygen levels were reported by [1975Fuk] (4 mass% O), [1981Ely] (0.15 mass%

B2O3), [1988Lor2] (sintered hBN, 0.6 mass% O, GI = 1.6), [1993Boc] (various grades of sintered and

powdered hBN with grain sizes from 0.5 to 40 m and 1.71 < GI < 2.08, mainly used 0.2 to 2.5 mass% O;

also used 4-6 mass% O, but received low yields of cBN), [1994Gla] (hBN powder, purified by a

magnesiothermic method, 0.6 mass% O).

In the following subsections we attempt to summarize in detail the experimental conditions only for all

those key-papers on which the hitherto accepted constitution of the pT diagrams is primarily based:

Mg-Mg3BN3 [1979End2] and Mg3BN3-BN [1979End2, 1993Nak1, 1995Lor1, 1995Lor2, 1997Lor].

Several review articles deal with the high pressure-high temperature cubic boron nitride synthesis in general

[1984Fuk, 1995Dem1, 1995Dem2] involving also supercritical fluids such as NH3 or NH2NH2 for

reduction of the cBN formation to 1.5 < p < 2.5 GPa and 500 < T < 700°C.

For studies of the influence of oxygen on the hBN to BN conversion process see section Miscellaneous.

Routes to reduce the oxygen level in hBN were presented by [1981Sat].

Binary Systems

A tentative B-Mg equilibrium diagram is given in [Mas2], revealing under normal conditions the formation

of three stable magnesium borides, MgB2, MgB4 and MgB7. Practically no mutual solid solubility exists

between B and Mg. Severe doubts concern the formation of “Mg3B2”, “MgB6” and “MgB12”. The crystal

structures of the magnesium borides MgB2 and orthorhombic “MgB6” (probably MgB7?), obtained from

the pressure and temperature range from 1223 to 1923°C and from 4 to 7 GPa, were said to remain

unchanged with respect to normal conditions [1967Fil]. MgB2 does not reveal any structural transitions up

to 40 GPa [2001Bor]. History of the B-Mg phase diagram has been summarized by [2003Lee] particularly

in relation to the problems in MgB2 crystal growth.

The partial phase diagram available in [Mas2] for Mg-N lists one compound, Mg3N2, with only one

modification. The phase diagram of Mg3N2 was investigated by [1993Gla] in the range from 1.5 to 9.0 GPa

and up to 1600°C and was said to contain a total of six solid phases of which three solid phases are present

at normal pressure. A kinetically slow transformation to a tetragonal high pressure phase (no structural

details revealed) was observed at 5.5 GPa and 897°C [1993Lor]. At 5.5 GPa the structure of hP-Mg3N2 was

said to persist up to the experimental temperature limit of 1570°C. Release of temperature and pressure

triggered a second transformation into a further phase with unknown structure and rather slow

transformation kinetics [1993Lor]. Interstitial solubility of N in Mg under normal conditions is very low.

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Phase relations for the B-N diagram have been presented by [2003Rec]; data on the p-T diagram have been

summarized by [2001Rog]. They have been accepted for this evaluation.

Solid Phases

Four ternary compounds have been found to exist in the ternary system B-Mg-N under various conditions

of temperature and pressure: Mg3B2N4, Mg3BN3, Mg6BN5 and MgB9N (for crystallographic data see

Table 1).

Much of the existing confusion among ternary magnesium boron nitrides was solved when the proper

stoichiometry of the ternary phase, first observed and labelled as “Mg3B2N4” by [1972Vri, 1979End1,

1979End2], was derived from X-ray Rietveld analyses to be Mg3BN3 [1991Hir, 1992Hir, 1993Hir,

1994Zhu, 1999Kul2]. It should be noted, that already [1986Sat] and [1989Pik] assigned the proper Mg3BN3

stoichiometry to this phase. The findings of [1986Sat] were based on DTA, DTG, electron and X-ray

powder diffraction data, whilst those of [1989Pik] refer to combined chemical and XRD-analyses. It is now

widely accepted, that Mg3BN3 is the only compound stable in the BN-Mg3N2 section under normal

conditions [1992Nak2, 1993Nak1, 1993Nak2, 1993Lor, 1994Zhu, 1995Lor1, 1995Lor2, 1997Lor].

Mg3BN3 exists in a low-pressure and a high-pressure modification: the p-T diagram of Mg3BN3 was studied

by [1993Nak1] and was essentially confirmed by [1994Zhu]. Figure 1 compares the findings of both

research groups, [1993Nak1] and [1994Zhu]. [1993Nak1] prepared lP-Mg3BN3 from Mg-flakes mixed

with hBN (3Mg+hBN) in a dry nitrogen atmosphere. After heating to 600°C and 650°C for 4h each, a first

reaction product was hBN+Mg3N2, which was crushed and reheated to 1200°C for 12 hrs under N2 in a

tightly capped steel capsule. At that stage the sample contained yellow lP-Mg3BN3 besides small amounts

of hBN, MgO and Mg(OH)2. Reaction at higher temperature causes decomposition of Mg3BN3 but lower

temperatures tend to reveal incompletely reacted hBN. The high pressure-high temperature form of

Mg3BN3 was prepared in a cubic 10 mm anvil high pressure apparatus. A sample disk was made compacting

lP-Mg3BN3 powder under dry nitrogen. The sample, encapsulated in Pt was then packed in a NaCl pressure

medium, pressurized to 4.0 GPa, after which temperature from an outer C-heater was increased to 1300°C

at a speed of 200°C min-1 and held at the set temperature for 15 min. After quenching the pressure was

slowly released. For details of pressure and temperature calibration see the original paper [1993Nak1].

Under a pressure of 0.1 MPa of N2, hP-Mg3BN3 decomposed and evaporated above 1150°C. At room

temperature lP-Mg3BN3 is stable below at least 5.8 GPa, but with increasing temperature transforms to the

high temperature high pressure form. The phase boundary was given as p(GPa) =7.2-0.0035T (°C) and

dp/dT (GPa/°C)=-3.5 10-3 [1993Nak1], whilst a monotonous but nonlinear transition curve at slightly lower

pressures was derived by [1994Zhu] (for comparison see Fig. 1). Infrared spectra of Mg3BN3 (“Mg3B2N4”)

formed at 2.5 GPa confirmed the linear [N=B=N]3- molecular ions [1979End1]. Mg3BN3 (“Mg3B2N4”)

slowly decomposes in moist environments; at higher temperatures, it easily oxidizes to Mg3(BO3)2

[1979End1].

With the determination of the crystal structures of the two forms of Mg3BN3 part of the confusion inherent

to earlier reports seems to be solved: “Mg3B2N4” as reported by [1979End1, 1979End2] is now established

as Mg3BN3; similarly “Mg3B2N4” of [1989Hoh] rather should be hP-Mg3BN3. Also “ Mg3B2N4” of

[1981Ely] in fact corresponds to hP-Mg3BN3. Whilst [1981Ely] in addition reported a decomposition of

their “ Mg3B2N4” into Mg3N2+cBN at 6 GPa and 800°C, neither decomposition of lP-Mg3BN3 nor of

hP-Mg3BN3 was detected by [1992Nak2, 1994Zhu] under dry conditions. With respect to these changes the

original description of the partial systems Mg-“Mg3B2N4” and “Mg3B2N4”-BN should rather read as

Mg-Mg3BN3 and Mg3BN3-BN. According to Endo [1979End1, 1979End2] Mg3BN3 melts under 2.5 GPa

at 1489 5°C and forms a eutectic with Mg at 737 1°C as well as with hBN at 1295 7°C. At this stage it

also should be noted, that the phase “Mg3BN3”, as labeled by Hohlfeld, [1989Hoh], was recognized by

Nakano, [1993Nak1] to correspond to his X phase, Mg6BN5 (see below). Accordingly, the melting point

reported by [1989Hoh] for “Mg3BN3” at 1685 10°C under 6.5 GPa rather applies to Mg6BN5 and so will

be the corresponding eutectic L Mg6BN5 (earlier “hP-Mg3BN3”)+BN at 6.5 GPa and 1380 10°C

[1989Hoh]. True Mg3BN3 is not the only compound in the pT diagram of B-Mg-N. [1986Yam] noticed the

appearance of two prominent X-ray intensities at d = 210 and 150 pm in Mg3N2 doped hBN after reaction

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B–Mg–N

at 5.8 GPa above 1175°C. This phase was tentatively labelled X(Chi)-phase [1986Yam]. From a detailed

investigation of the BN-Mg3N2 region (2<(BN/Mg3N2)<0.1) in the range from 2 to 5.5 GPa and from 1000

to 1500°C the authors of [1993Nak1] confirmed the existence of the X phase, although they were unable to

obtain a single phase X-ray diffraction pattern. As a result of their investigation they concluded, that no

high-pressure compounds were found in the subsystem BN-Mg3BN3 and no decomposition reaction

forming cBN was observed. From samples in the region Mg3N2-Mg3BN3, reacted for 15 min at 4.0 to 4.4

GPa and 1300°C, it was concluded, that the composition of the X phase might be slightly richer in Mg3N2

than Mg3BN3 consequently a tentative formula “Mg6BN5” was introduced. There was neither evidence of

decomposition of hP-Mg3BN3 into the X phase and BN, nor evidence of decomposition of the X phase to

form cBN. Adding hBN to the sample containing the X phase reduced its quantity and above 1300°C the X

phase was said to become unstable. Comparing the X-ray data presented, [1993Nak1] found that “hP-

Mg3BN3” of [1989Hoh] in fact corresponds to the X phase. A phase labeled Mg6BN5 was indeed observed

as a brick-red powder by [1999Kul1, 1999Kul2, 1999Kul3, 2000Kul] in high-pressure - high-temperature

experiments using a toroid-type equipment and starting from various mixtures MxNy-BN. The phase was

observed in the region from 1.5 to 5.0 GPa near 1100°C with significant amounts of secondary phases,

mostly hBN and Mg3BN3; the best sample, almost free of Mg3BN3 but with small amounts of h-BN and

MgO, was found at an optimal pressure of 2.0 GPa from an initial mixture 2 Mg3N2+hBN and 1600°C for

1h [1999Kul1, 2000Kul]. At lower pressure formation of lP-Mg3BN3 occurred and at increased pressure

hP-Mg3BN3 appeared [1999Kul1, 1999Kul2, 2000Kul]. At 1100°C [1999Kul2] reported a critical pressure

of 4.5 GPa above which formation of hP-Mg3BN3 was observed. The X-ray powder pattern of Mg6BN5 was

tentatively indexed on the basis of a hexagonal unit cell (a = 539.7 pm, c = 1058.5 pm); the crystal structure

has not been elucidated [1999Kul1]. Mg6BN5 was said to gradually react with moisture [1999Kul1]. As

Mg6BN5 was found to be unstable above 5.0 GPa, its role in h-BN to c-BN conversion is limited to the

formation of hP-Mg3BN3 according to the reaction: Mg6BN5+hBN 2hP-Mg3BN3 [1999Kul1]. Based on

infrared spectra, the structure of Mg6BN5 was said to contain N3- and (N=B=N)3-anions similar to Mg3BN3

[1999Kul3, 2000Kul].

From in-situ synchrotron experiments in a multi-anvil press at 5.5 GPa on powder blends of 5Mg3N3+hBN

the authors of [1995Lor1] observed the formation of a further compound above 1277°C, which was

concluded to be richer in BN than Mg3BN3. This new compound (with a composition close to Mg3B2N4?)

was claimed to be the solvent for the cBN formation. No reaction was found on quenching down to 4.0 GPa

and 200°C suggesting metastability for the ternary compound. But at 0.5 GPa, 200°C cBN formation was

observed with decomposition of the ternary compound and eventually formation of a further ternary phase

of yet unknown crystal structure. All these results were used to construct a phase diagram Mg3BN3-BN (see

below). High pressure experiments in the range from 1050 to 1600°C, 4 to 6 GPa for 1 hour on various

mixtures of Mg3N2+hBN or Mg3BN3+hBN were found to yield a green powder of a new phase, which was

tentatively indexed on the basis of a hexagonal system (a = 1339.45, c = 595.17 pm) and which was labeled

as Mg3B2N4 [1999Kul2]. After annealing at 900°C under nitrogen this phase was said to decompose into

lP-Mg3BN3 [1999Kul2]. The authors of [1999Kul2] assumed that cBN formation takes place above 4.5 GPa

via formation of hP-Mg3BN3 and Mg3B2N4 according to the reaction:

Mg3BN3+hBN Mg3B2N4 hP-Mg3BN3+cBN; Mg3BN3 is either directly obtained from Mg3N2+hBN or

from 2Mg3N2+hBN Mg6BN5+hBN 2Mg3BN3.

The novel boron framework structure of MgB9N, consisting of boron icosahedral linked to boron

octahedral, was characterized from single crystal data; the structure may alternatively be described as an

intergrowth of NB6 and MgB3 layers [2002Mir]. MgB9N was synthesized from a 5Mg-1B mixture in a

BN-crucible in a high gas-pressure apparatus, applying first argon pressure (100 MPa) and then raising the

temperature to 1600°C, holding it there for 1 h and then cooling from 1600°C to 1500°C at a rate of 60°C/h.

Below 1500°C the sample was cooled to room temperature art a rate of 600°C/h. To remove excess Mg, the

product was then heated in vacuo at 750°C for 15 min, yielding black single crystals analyzed by EMPA to

give the formula MgB9N.

It is furthermore interesting to note, that the authors of [1990Fut] and [1992Evd] described an unknown

phase, labeled X, in the system MgB2-BN, supposed to be involved in the production of cBN from a eutectic

melt.

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B–Mg–N

Isothermal Sections

Figure 2 is a schematic presentation of the phase relations in the B-Mg-N system at moderate temperatures

(about 400°C) extrapolated from the findings available in literature.

The pT diagrams Mg-BN, Mg-Mg3BN3, MgB2-BN and Mg3N2-BN.

Phase relations along the Mg-Mg3N2-BN boundaries were determined at 2.5 GPa in the temperature range

from 800 to 1600°C employing DTA, optical, and X-ray powder diffraction analysis supported by

conventional quenching techniques [1979End2]. The experiments were done in pyrophyllite/BN or

tantalum cells in a belt apparatus where pressure was applied to a given value and the temperature was raised

at a rate of 200 to 250°C min-1. Temperature and pressure were held for 20 min followed by quenching to

room temperature within ~30 sec with the pressure still applied. No reaction was found between Mg and

BN below 1100°C, whereas above 1150°C yellow lath-like plates of ternary magnesium boron nitride,

labelled as “Mg3B2N4” (now Mg3BN3) and magnesium borides “MgB~6”, “MgB~12” were found. The

precipitation of translucent cBN with high thermal conductivity from the eutectic liquid was observed for

temperatures above 1300°C [1979End2].

In a cursory investigation of the p-T diagram of the Mg3N2-BN system for compositions Mg3N2+nBN

(0.5<n<4.0) [1981Ely] claimed the ranges of existence of two modifications of a ternary compound at 4.0

GPa, which from chemical analysis were specified as “ Mg3B2N4” (800 to 1200°C), “ Mg3B2N4”

(>>1300°C) with an intermediate two-phase region in between. The minimum pressure necessary for cBN

formation at 800°C was reported as 4.5 GPa, raising to 6 GPa at 800 to 1100°C. A study of the kinetics of

the cBN-synthesis from the three-phase fields showed (after an exposure time of 10 min to 6 and 7 GPa) a

steady increase of the cBN-yield from 900°C to about 1280°C (80%) followed by a pronounced minimum

(40%) at 1300°C reassuring high yields above 1423°C. It was assumed, that the formation of “ Mg3B2N4”

leads to the decrease of the reaction rate and a decrease in the yield of cBN. On further heating “ Mg3B2N4”

becomes unstable and decomposes. The X-ray powder pattern as reported by [1979End1] for “Mg3B2N4”

was said to correspond essentially to this two-phase region consisting of “ Mg3B2N4”+“ Mg3B2N4”

[1981Ely]. The existence of two modifications of “Mg3B2N4” was not confirmed by [1989Hoh], who

claimed coexistence of his “Mg3B2N4” with a second compound “Mg3BN3” within the low-temperature

cBN-region [1989Pik, 1989Hoh]. “Mg3BN3” was assumed to act as a chemical catalyst in the rapid

solid-state reaction “Mg3BN3”+hexagonal BN “Mg3B2N4” “Mg3BN3”+cBN at 1330°C for 6.5 GPa

[1989Hoh]. Here it should be stressed again, that Hohlfeld’s phase “Mg3BN3”, in fact corresponds to the X

phase, Mg6BN5, and that Hohlfeld’s phase “Mg3B2N4” rather should be hP-Mg3BN3. With these

substitutions, however, the reaction given by Hohlfeld [1989Hoh] will not be satisfied.

The investigation of the high-pressure reactions in the section Mg3N2-hBN on a series of samples reacted

at 4.0 GPa and 1300°C for 15 min revealed a new high-pressure phase (X phase) slightly richer in Mg3N2

than Mg3BN3 [1993Nak1]. At 1000°C only a small amount of this phase was produced, the amount

increasing with pressure. There was no evidence for a decomposition of hP-Mg3BN3 into the X phase+BN.

It shall be noted, that according to the conclusions of [1993Nak1], [1989Hoh] mis-assigned the hP-Mg3BN3

phase for the X phase. With additions of BN the X phase reacts to form hP-Mg3BN3 but at 1300°C there is

competition to form cBN. 4.9 GPa was established as the critical minimum pressure to form cBN from hBN.

There was, however, no evidence of a decomposition of the X phase to form cBN. It was confirmed that

cBN is produced from the reaction hP-Mg3BN3 and hBN only above 1300°C, which is close to the melting

point of hP-Mg3BN3 [1993Nak1].

A high pressure-high temperature kinetic and thermodynamic investigation of cBN formation in the system

BN-Mg3N2 is due to [1995Lor1, 1995Lor2, 1997Lor] employing time resolved in-situ synchrotron

diffraction and a specially designed multi-anvil device. As a result of this investigation a partial section of

the Mg3N2-BN phase diagram was constructed at 5.5 GPa (Fig. 4). The phase diagram confirms the

formation of Mg3BN3 but also states the peritectic formation of a new compound slightly richer in BN than

Mg3BN3. The composition of this new phase was tentatively assumed by [1997Lor] as Mg3B2N4 but with

a crystal structure different from the phase earlier designed as “Mg3B2N4” by [1979End2, 1981Ely,

1989Hoh]. At 5.5 GPa Mg3BN3 is engaged in a eutectic with the “new” Mg3B2N4, which forms

peritectically: L+BN Mg3B2N4 at a temperature slightly below 1417°C. It was not clearly reported if and

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B–Mg–N

under what conditions the new phase decomposes, although in the paper by [1995Lor1] the decomposition

temperature of Mg3B2N4 was shown to coincide with the eutectic reaction L Mg3B2N4+Mg3BN3!

When samples of Mg3N2+BN were heated a metastable eutectic L Mg3BN3+BN was observed at 1227°C

followed by a metastable peritectic reaction L+BN Mg3B2N4 at 1417°C. At 5.5 GPa the transition

hBN cBN was analyzed at 1223°C and 1347°C with respect to kinetic models.

Being aware of the persisting confusion in the description of the phase relations we show the phase diagram

version of [1979End2] at 2.5 GPa (corrected by a shift of the composition from “Mg3B2N4” to Mg3BN3)

(subsystems Mg-Mg3BN3 and Mg3BN3-BN) in Fig. 3 and the version of [1997Lor] at 5.5 GPa in Fig. 4

(subsystem Mg3BN3-BN).

A treshold pressure of 4.5(1) GPa and low temperature boundary of T(K) = 1633-9.2p (GPa) was

established by [1999Sol] for the formation of cBN crystallizing from BN-MgB2 melts. The experiments

were performed in a WC-multi anvil press and synchrotron radiation. Below 4.4 GPa no formation of cBN

has been detected up to 1800°C. The existence of a ternary eutectic L MgB4+Mg3BN3+BN was suggested.

Accordingly the section MgB2-BN is not a quasibinary system [1999Sol]. For successful growth of MgB2

single crystals in BN containers [2003Lee] see section Miscellaneous.


According to the descriptions of the phase relations in the two subsystems investigated, Mg-Mg3BN3 and

Mg3BN3-BN, the eutectic reaction L Mg3BN3+hBN (hexagonal) at 1295 7°C at 2.5 GPa [1979End2]

corresponds to the metastable reaction of [1997Lor] assigned at about 1227°C and 5.5 GPa.

The pressure dependence of the metastable reactions hBN+Mg3N2, hBN+Mg3BN3 and hBN+“Mg3B2N4”

was measured by in-situ DTA in a toroid-type apparatus in the pressure range from 3.0 to 8.0 GPa and for

temperatures up to 1627°C (using Pt/Pt-Rh thermocouples) [1994Gla]. On various starting mixtures with

ratio of hBN: Mg3N2 (containing 0.6 mass% O) varying from 1:3 to 3:1, the pressure was first increased

and then the temperature (1.2°CMg s-1). The authors of [1994Gla] claim a eutectic nature for all the three

reactions mentioned above (see Fig. 5), however, a detailed in-situ investigation by synchrotron radiation

of the high pressure-high temperature formation of Mg3BN3 [1993Lor] suggests a solid state reaction of

hBN+Mg3N2 Mg3BN3 at 887 20°C. The reaction was observed without any melting features and was

said to be nearly pressure independent in the investigated range from 4.3 to 5.5 GPa. Correspondence with

data by [1994Gla] was far from excellent but was said to be acceptable [1993Lor]. The reaction

Mg3BN3+hBN essentially corresponds to the metastable eutectic L Mg3BN3+hBN at 1227°C at 5.5 GPa

given by [1997Lor] in Fig. 4. According to [1997Lor] the reaction Mg3B2N4+hBN seems to rather reveal

the peritectic formation L+hBN Mg3B2N4 of the new phase Mg3B2N4 at 1417°C at 5.5 GPa (Fig. 4).

Miscellaneous

The phase transformation of amorphous boron nitride (aBN) to cBN in the presence of magnesium boron

nitride was studied for the ratio aBN: “Mg3B2N4” (Mg3BN3?) = 1:1 in a belt type high pressure apparatus

for the pT range 2.0 < p < 7.0 GPa and 727 < T < 1827°C [1995Sin1, 1995Sin2]. The minimum pressure of

cBN formation was 2.5 GPa at 1800°C raising to 4 GPa at 1200°C and to 7 GPa at 900°C [1995Sin2]. Whilst

cBN formation at temperatures higher than 1300°C and pressures greater than 4.0 GPa was conceived via

a catalyst-solvent process through a B-Mg-N melt, cBN crystallization at temperatures as low as 900°C and

7 GPa were rather explained in a direct transformation mechanism [1995Sin2]. At temperatures higher than

1300°C and 4.0 GPa respectively, formation of MgB4 was observed, which was interpreted as a dissociation

product of the magnesium boron nitride catalyst [1995Sin2].

Additions of boron (10 to 25 mass%) or Al (25 mass%) to the Mg3N2 solvent assisted cBN production from

hBN at 1700°C and at a minimum pressure of 5 GPa were said to increase the yield of cBN by a factor of

about 5 with respect to pure Mg but without significant improvement of morphology and translucency of

the cBN crystals [1993Boc].

The influence of phosphorus on the crystallization kinetics of cBN in the MgB2-BN system was studied by

[1989Shi] at 5 GPa and 1570 to 1800°C; 0.1 to 0.5 mass% phosphorous considerably increases the

nucleation rate and reduces the nucleation energy for sphalerite-type BN.

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B–Mg–N

The role of oxygen in the Mg or Mg3N2 -assisted hBN-cBN formation.

Most experiments suffer from small amounts of oxygen and moisture introduced via the highly hygroscopic

Mg3N2 starting material or by B2O3, a common impurity to hexagonal boron nitride (hBN). According to

the reaction, B2O3+Mg3N2 3MgO+2BN, various amounts of MgO are obtained in the final products.

Similar to Mg3N2, the ternary magnesium boron nitrides tend to hydrolyze in moist air yielding Mg(OH)2

and ammonia.

The role of oxygen in the hBN-cBN conversion and its influence on the minimum pT conditions as well as

on the cBN yields was the subject of a series of investigations with some contradictory results.

Differentiating between the Mg-BN and the Mg3N2 systems, Figs. 6 and 7 summarize the cBN growth

regions determined by the various groups as a function of their oxygen levels in the hBN starting material.

Although for the Mg-BN system there seems to be a distinct separation into two sets of curves interpretation

is not straight forward: whereas [1979End1, 1993Boc] observed a remarkable decrease of minimum

pressure and temperature necessary for cBN formation for low oxygen contents, the data of [1991Bin1,

1991Bin2] for their high oxygen hBN (9 mass% O) rather correspond to the low oxygen data of [1979End1,

1993Boc] (2 and 0.2 mass% O). [1979End1] concluded that the growth of cBN was significantly obstructed

by the formation of oxides MgO and/or Mg3(BO3)2, consuming the catalyst/solvent. Therefore the high

levels in MgO by-product of [1966Kud] group their data into the high oxygen hBN materials among the

second set of curves with higher minimum pT conditions for the cBN formation. An explanation for the

discrepancy with the data of [1991Bin1, 1991Bin2] was offered by [1993Boc], who compared the different

experimental set-up of [1991Bin1, 1991Bin2] (hBN-Mg powder blends) and of [1979End1, 1993Boc]

(alternately piled discs of hBN-Mg-hBN). Following the argumentation of [1993Boc], the powder mixture

allows a rapid formation of large numbers of cBN nuclei before the retardation by MgO sets in, whilst in

the stacked pile MgO accumulates at the interface and efficiently hinders further nucleation and growth of

cBN.

[1981Sat] extended the investigation to the system of Mg3N2-BN demonstrating that oxygen has little effect

on the lower pressure and temperature limits of the cBN formation (see also Fig. 6). However, high cBN

yields only resulted from low oxygen hBN. High oxygen levels in hBN starting material result in the

formation of MgO and abnormal cBN crystal morphologies (see also [1988Lor1] below). Efficient

reduction in oxygen content was achieved with Mg3N2+Zr getter-mixtures increasing the yield of cBN

crystals with smooth crystal surfaces.

The influence of MgO containing catalyst/solvent Mg3N2 on the growth of cBN was studied from powder

blends 5Mg3N2+MgO (in mass%) reacted in a belt-type high pressure cell at 1527°C and 6.5 GPa. A 2.5 to

5 m thick layer of MgO formed around the cBN grains which easily dissolved in diluted HCl leaving small

pits on the cBN crystal surfaces [1988Lor1]. In contrast to the findings of [1981Sat] no anomalous cBN

crystal morphologies were observed.

The important role of Mg3BN3 in the industrial production of translucent cBN has to be stressed [1984Sat].

Additions of small amounts of H2O to Mg3BN3-containing hBN were said to lead to better quality

translucent cBN [1991Sum, 1992Nak2].

Alternative flux/solvent systems containing Mg.

Considerable efforts were made to develop alternative flux solvents for the hBN to cBN conversion. A

summary is given on those materials which contain magnesium as magnesium fluoride (MgF2) [1979Kob],

magnesium fluoronitrides (Mg2NF, Mg3NF3) [1986Dem, 1995Dem1] or mixed flux systems

(“Mg3B2N4”+xLiF) [1992Dem, 1995Dem1]. Spontaneous crystallization of sphaleritic BN was observed

from mixtures of 60% hBN+10% MgB2+30mass% NH3 after 1000 sec at 1027 to 1227°C and at pressures

higher than 2.1 0.1 GPa [1992Sol1, 1994Sol]. A similar range of conditions was reported for a starting

mixture of hBN+Mg3N2+NH3 [1992Sol2, 1994Sol]. For these experiments a low-friction cell for a

piston-cylinder toroid-type high pressure apparatus was used and fresh hexaammincalcium(II)chloride as

source of ammonia [1992Sol1, 1992Sol2]. In the presence of supercritical NH3 the threshold pressure for

spontaneous crystallization of cBN on cBN seed crystals is reduced at 1027°C as low as 0.5 GPa [1994Sol].

103


MSIT®

B–Mg–N

High quality 200-300 m single MgB2 crystals of different morphology were grown from Mg+B powder

mixtures in BN containers at 5 GPa, 1600-1700°C, with a temperature gradient of >20°C/mm in a liquid

assisted solid state re crystallization mechanism [2002Lee, 2003Lee]. They demonstrated a sharp

superconducting transition at Tc = 38.1 - 38.3 K in both magnetization ( Tc=0.6 K) and resistivity

( Tc=0.3 K) measurements.

References

[1960Wen] Wentorf, R.H.Jr., “Abrasive Material and Preparation Thereof”, U.S. Patent 2947617 GEC,

2 Aug (1960)

[1966Kud] Kudaka, K., Konno, H., Matoba, T., Kogyo Kagaku Zasshi (J.Chem. Soc. Jpn., Ind. Chem.

Section), 69, 365 (1966) (Experimental, 4)

[1967Fil] Filonenko, N.E., Ivanov, V.I., Feld’gun, L.I., Sokhor, M.I., Vereshchagin, L.F.,

“Magnesium Borides Prepared Under Superhigh Pressure Conditions”, Dokl. Akad. Nauk

SSSR, 175, 1266-1269 (1967) translated from Sov. Phys. Dokl., 12, 833-835 (1968)

(Experimental, Crys. Structure, 5)

[1971Ush] Ushio, M., Saito, H., Nagano, S., Kogyo Kagaku Zasshi (J. Chem. Soc. Jpn.,

Ind.Chem.Section), 74, 598 (1971) (Experimental, 6)

[1972Vri] de Vries, R.C., Fleischer, J.F., “Phase Equilibria Pertinent to the Growth of Cubic Boron

Nitride”, J.Cryst. Growth, 13/14, 88-92 (1972) (Experimental, Equi. Diagram, 9)

[1975Fuk] Fukunaga, O., Sato, T., Iwata, M., Hiraoka, H., “Synthesis of cBN by the Catalyst Process”,

Proc. 4th Intl. Conf. on High Pressure, Kyoto (1974), Osugi, J., Suzuki, K., Koizumi, M.,

Makita, T., Matsushima, S., (Eds.), Phys. Soc. Japan, Tokyo, 454-59 (1975)

(Experimental, 10)

[1979Kob] Kobayashi, T., “Solvent Effects of Fluorides in cBN High Pressure Syntheses”, Mater. Res.

Bull., 14, 1541 (1979) (Experimental, Equi. Diagram, 7)

[1979End1] Endo, T., Fukunaga, O., Iwata, M., “Growth Pressure-Temperature Region of Cubic BN in

the System BN-Mg”, J. Mater. Sci., 14, 1375-1380 (1979) (Experimental, Diagram, 15)

[1979End2] Endo, T., Fukunaga, O., Iwata, M., “Precipitation Mechanism of BN in the Ternary System

of B-Mg-N”, J. Mater.Sci. 14, 1676-1680 (1979) (Experimental, Crys. Structure, Equi.

Diagram, 8)

[1981Ely] Elyutin, V.P., Polushin, N.I., Burdina, K.P., Polyakov, V.P., Kalashnikov, Y.A.,

Semenenko, K.N., Pavlov, Yu.A., “Interaction in the System Mg3N2-BN”, Dokl. Akad.

Nauk SSSR, 259, 112-116 (1981) (Experimental, Crys. Structure, Diagram, 6)

[1981Sat] Sato, T., Hiraoaka, H., Endo, T., Fukunaga, O., Iwata, M., “Effect of Oxygen on the Growth

of Cubic Boron Nitride using Mg3N2 as Catalyst”, J. Mater. Sci., 16, 1829-1834 (1981)

(Experimental, Crys. Structure, Equi. Diagram, 6)

[1984Fuk] Fukunaga, O., “High Pressure Synthesis of Hard Materials: C, BN”, J. Phys., Colloq., C8

(11), 45, 315-324 (1984) (Experimental, Review, 19)

[1984Sat] Sato, T., Endo, T., Fukunaga, O., US Patent 532093, (1984)

[1986Dem] Demazeau, G., Biardeau, G., Pouchard, M., Brevet Fr. 2.597.087 (1986)

[1986Sat] Sato, T., NIRIM Research Report No.46 (in Japanese), 53-58 (1986) (Experimental, Crys.

Structure, 8)

[1986Yam] Yamaoka, S., Shimomura, O., Akaishi, M., Kanda, H., Nagashima, T., Fukunaga, O.,

Akimoto, S., “X-Ray Observation of the Formation of Diamond and Cubic Boron Nitride at

High Pressure and Temperature”, Physica, 139&140B, 668-670 (1986) (Experimental,

Crys. Structure, Equi. Diagram, 3)

[1988Lor1] Lorenz, H., Kühne, U., Hohlfeld, C., Flegel, K., “Influence of MgO on the Growth of Cubic

Boron Nitride Using the Catalyst Mg3N2”, J. Mater. Sci. Lett., 7, 23-4 (1988)

(Experimental, 5)

[1988Lor2] Lorenz, H., Lorenz, B., Kühne, U., Hohlfeld, C., “The Kinetics of Cubic Boron Nitride

Formation in the Sytem BN - Mg3N2”, J. Mater. Sci., 23, 3254-7 (1988) (Experimental, 15)

104


MSIT®

B–Mg–N

[1989Hoh] Hohlfeld, C., “Peculiarities of the Cubic Boron Nitride Formation in the System

BN-Mg3N2”, J. Mater. Sci. Lett., 8, 1082-1084 (1989) (Experimental, Crys. Structure, 6)

[1989Pik] Pikalov, S.N., Tarisenko, N.V., Sobokary, O.A., Bondarevich, K.A., “Composition of

Magnesium Boronitride Synthesized from Boron Nitride and Magnesium in a Nitrogen

Atmosphere at 1400 K”, Poroshk. Metall., 28(3) 71-74 (1989) (Experimental, 8)

[1989Shi] Shipilo, V.B., Gameza, L.M., “Crystallization of Sphalerite-Type Boron Nitride in the

BN-MgB2-P System”, translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 25, 51-55

(1989) (Experimental, 8)

[1990Fut] Futerhendler, S.I., Davydenko, V.M., Shemakin, V.I., Superhard Materials, (3), 22-24

(1990) (Experimental)

[1991Bin1] Bindal, M.M., Singh, B.P., Singhal, S.K., Nayar, R.K., Chopra, R., Dhar, A., “On the

Choice of Hexagonal Boron Nitride for High Pressure Phase Transformation Using the

Catalyst Solvent Process”, J. Mater. Sci., 26, 196-202 (1991) (Experimental, Crys.

Structure,18)

[1991Bin2] Bindal, M.M., Singhal, S.K., Singh, B.P., Nayar, R.K., Chopra, R., Dhar, A., “Synthesis of

cBN Using Magnesium as the Catalyst”, J. Cryst. Growth, 112, 386-401 (1991)


[1991Hir] Hiraguchi, H., Hashizume, H., Fukunaga, O., Takenaka, A., Sakata, M., “Structure

Determination of Mg3BN3 from X-ray Powder Diffraction Data”, J. Appl. Crystallogr., 24,

286-292 (1991) (Experimental, Crys. Structure, 43)

[1991Sum] Sumiya, H., Tsuji, K., Yazu, S., in “New Diamond Science and Technology”, Messier, R.,

Glass, J.T., Butler, J.E., Roy, R., (Eds.), Mater. Res. Soc., Pittsburgh, PA, p.1063 (1991)

(Experimental)

[1992Dem] Demazeau, G., “New Routes for Preparing Cubic Boron Nitride: Synthesis and Crystal

Growth”, in Recent Trends in High Pressure Research, Proc. XIII AIRAPT Intl. Conf. on

High Pressure Science and Technology, 1991, Singh, A.K. (Ed.), Oxford IBH

Publ.Co.PVT.LTD., New Delhi (1992) p.681-686 (Review, 60)

[1992Evd] Evdokimova, O.M., Pesin, V.A., Borovikov, S.V., Davydenko, V.M., in “Physical

Chemistry and Processes of Obtaining High Quality Abrasive Materials by Ecologically

Secure Techniques”, St. Petersburg, VNIIAH, (1992), 77-81 (Review)

[1992Hir] Hiraguchi, H., Hashizume, H., Nakano, S., Fukunaga, O., Abstracts of Accuracy in Powder

Diffraction II, Hawai, 1992, p.41

[1992Lor] Lorenz, H., Orgzall, I., Hinze, E. Kremmler, J., “Structural Phase Transitions of Mg3BN3

under high Pressures and Temperatures”, Scr. Metall. Mater., 287, 993-997 (1992)


[1992Nak1] Nakano, S., “New Concept of the Synthesis of Cubic Boron Nitride”, in Recent Trends in

High Pressure Research, Proc. XIII AIRAPT Intl. Conf. on High Pressure Science and

Technology, 1991, Singh, A.K. (Ed.), Oxford IBH Publ.Co. PVT.LTD., New Delhi (1992)

p.687-691 (Experimental, Equi. Diagram, 14)

[1992Nak2] Nakano, S., Ikawa, H., Fukunaga, O., “Synthesis of Cubic Boron Nitride by Decomposition

of Magnesium Boron Nitride”, J. Am. Ceram. Soc., 75(1), 240-43 (1992) (Experimental,


[1992Sol1] Solozhenko, V.L., Slutskii, A.B., Ignatiev, Y.A., “On the Lowest Pressure of Sphaleritic

Boron Nitride Spontaneous Crystallization”, J. Superhard Materials, 14, 64 (1992)

(Experimental, 3)

[1992Sol2] Solozhenko, V.L., “A New Boron Nitride Phase Diagram and Chemical Interaction in the

B-N System at High Pressures”, High Pressure Res., 9, 140-143 (1992) (Experimental,

Equi. Diagram, 5)

[1993Boc] Bocquillon, G., Loriers-Susse, C., Loriers, J., “Synthesis of Cubic Boron Nitride Using Mg

and Pure or M’-Doped Li3N, Ca3N2 and Mg3N2 with M’=Al,B,Si,Ti”, J. Mater. Sci., 28,


105


MSIT®

B–Mg–N

[1993Hir] Hiraguchi, H., Hashizume, H., Sasaki, S., Nakano, S., Fukunaga, O., “Structure of a

High-Pressure Polymorph of Mg3BN3 Determined from X-Ray Powder Data”, Acta

Crystallogr., B49, 478-483 (1993) (Experimental, Crys. Structure, 18)

[1993Gla] Gladkaya, I.S., Kremkova, G.N., Bendeliani, N.A., “Phase Diagram of Magnesium Nitride

at High Pressures and Temperatures”, J. Mater. Sci. Lett., 12, 1547-1548 (1994)


[1993Lor] Lorenz, H., Orgzall, I., Hinze, E., Kremmler, J., “Formation of Mg3BN3 under High

Pressure and Temperature in the System Mg3N2-hBN”, Scr. Metall. Mater., 28, 925 (1993)


[1993Nak1] Nakano, S., Ikawa, H., Fukunaga, O., “High Pressure Reactions and Formation Mechanism

of Cubic Boron Nitride in the System BN-Mg3N2”, Diamond and Rel. Mater., 2, 1168-74

(1993) (Experimental, Crys. Structure, Equi. Diagram, 18)

[1993Nak2] Nakano, S., Ikawa, H., Fukunaga, O.,“Synthesis of Cubic Boron Nitride using Li3BN2,

Sr3B2N4 and Ca3B2N4 as Solvent-Catalysts”, Diamond and Rel. Mater., 3, 75-83 (1993)


[1993Sha] Shankland, K., Gilmore, C.J., Bricogne, G., Hashizume, H., “A Multisolution Method of

Phase Determination by Combined Maximisation of Entropy and Likelihood. VI.

Automatic Likelihood Analysis via the Student t Test, with an Application to the Powder

Structure of Magnesium Boron Nitride, Mg3BN3”, Acta Crystallogr., A49, 493-501 (1993)

(Crys. Structure, 30)


of Carbon-Doped -Rhombohedral Boron”, Phys. Status Solidi, B179, 489-511 (1993)


[1994Gla] Gladkaya, I.S., Kremkova, G.N., Bendeliani, N.A., Lorenz, H., Kühne, U., “The Binary

System of BN-Mg3N2 under High Pressures and Temperatures”, J. Mater. Sci., 29,

6616-6619 (1994) (Experimental, Crys. Structure, Equi. Diagram, 11)

[1994Sol] Solozhenko, V.L., “New Concept of BN Phase Diagram: An Applied Aspect”, Diamond

and Rel. Mater., 4, 1-4 (1994) (Experimental, Equi. Diagram, 40)

[1994Zhu] Zhukov, A.N., Burdina, K.P., Semenenko, K.N., “Composition and Polymorphism of

Magnesium Boron Nitride”, Russ. J. Gen. Chem., 64(3), Part 1, 323-326 (1994)

(Experimental, Equi. Diagram, Crys. Structure, 6)

[1995Dem1] Demazeau, G., “High Pressure Diamond and Cubic Boron Nitride Synthesis”, Diamond and

Rel. Mater., 4, 284-287 (1995) (Review, 50)

[1995Dem2] Demazeau, G., “Cubic BN: a New Route for Preparation at Ambient Pressures and

Temperatures” (in French), Compt. Rend. Acad. Sci. Paris, Ser. IIb, 320, 419-422 (1995)

(Review, Equi. Diagram, 16)

[1995Kur] Kurdyumov, A.V., Solozhenko, V.L., Zelyavsky, W.B., “Lattice Parameters of Boron

Nitride Polymorphous Modifications as a Function of Their Crystal-Structure Perfection”,

J. Appl. Crystallogr., 28, 540-545 (1995) (Crys. Structure, Experimental, 25)

[1995Lor1] Lorenz, H., Orgzall, I., “Formation of Cubic Boron Nitride in the System Mg3N2-BN: a

New Contribution to the Phase Diagram”, Diamond and Rel. Mater., 4, 1046-1049 (1995)


[1995Lor2] Lorenz, H., Orgzall, I., Hinze, E., “Rapid Formation of Cubic Boron Nitride in the System

Mg3N2-hBN”, Diamond and Rel. Mater., 4, 1050-1055 (1995) (Experimental, Crys.

Structure, 23)

[1995Sin1] Singh, B.P., Solozhenko, V.L., Will, G., “On the Low-Pressure Synthesis of Cubic Boron

Nitride”, Diamond and Rel. Mater., 4, 1193-5 (1995) (Experimental, Equi. Diagram, 13)

[1995Sin2] Singh, B.P., Nover, G., Will, G., “High Pressure Phase Transformation of Cubic Boron

Nitride from Amorphous Boron Nitride as the Catalyst”, J. Cryst. Grow., 152, 143-149

(1995) (Experimental, Equi. Diagram, Crys. Structure, 13)

106


MSIT®

B–Mg–N

[1996Hor] Horiuchi, S., He, L-L., Onoda, M., Akaishi, M., “Monoclinic Phase of Boron Nitride

Appearing During the Hexagonal to Cubic Phase Transition at High Pressure and High

Temperature”, Appl. Phys. Lett., 68(2), 182-184 (1996) (Experimental, Crys. Structure, 16)

[1997Lor] Lorenz, H., Peun, T., Orgzall, I., “Kinetic and Thermodynamic Investigation of cBN

Formation in the System BN-Mg3N2”, J. Appl. Phys., A65, 487-495 (1997) (Experimental,


[1999Kul1] Kulinich, S.A., Zhukov, A.N., Sevastyanova, L.G., Burdina, K.P., “On Some Alkali- and

Alkaline Earth-Metal Boron Nitrides Unsaturated with Boron”, Diamond and Rel. Mater.,

8, 2152-2158 (1999) (Experimental, Crys. Structure, 31)

[1999Kul2] Kulinich, S.A., Sevastyanova, L.G., Burdina, K.P., Semenenko, K.N., “Interaction in the

Mg3N2 - BN System under High Pressure”, Zh. Obshchey Khimii, 69(3), 358-363 (1999)


[1999Kul3] Kulinich, S.A., Sevastyanova, L.G., Bondarenko, G.N., Burdina, G.N., “On the Presence of

(N=B=N)3- Anions in some Boron Nitride Structures”, Zh. Obshchey Khimii, 69(4),

551-554 (1999) (Experimental, 14)

[1999Sol] Solozhenko, V.L., Turkevich, V.Z., Holzapfel, W.B., “On Nucleation of Cubic Boron

Nitride in the BN-MgB2 System”, J. Phys. Chem. B, 103, 8137-8140 (1999) (Experimental,

Equi. Diagram, 13)

[2000Kul] Kulinich, S.A., Sevastyanova, L.G., Zhukov, A.N., Burdina, K.P., “Boron Nitrides of Alkali

and Alkaline Earth Metals Containing N3- Anions”, Zh. Obshchey Khimii, 70(2), 190-196

(2000) (Experimental, Crys. Structure, 30)

[2001Bor] Bordet, P., Mezouar, M., Nunez-Regueiro, M., Monteverde, M., Nunez-Regueiro, M.D.,

Rogado, N., Regan, K.A., Hayward, M.A., He, T., Loureiro, S.M., Cava, R.J., “Absence of

a Structural Transition up to 40 GPa in MgB2 and the Relevance of Magnesium

Nonstoichiometry”, Phys. Rev. B, 64(17), 172502, pp 1-4 (2001) (Experimental, Crys.


[2001Pas] Paskowicz, W., Knapp, M., Domagala, J.Z., Kamler, G., Posiadlo, S.,“Low Temperature

Thermal Expansion of Mg3N2”, J. Alloys Compd., 328, 272-275 (2001) (Experimental,

Crys. Structure, 18)

[2001Rog] Rogl, P., “Materials Science of Ternary Metal Boron Nitrides”, Int. J. Inorg. Mater., 3,

201-209 (2001) (Review, Equi Diagram, Crys. Structure, 65)

[2002Oik] Oikawa, K., Kamiyama, T., Mochiku, T., Takeya, H., Furuyama, M., Kamisawa, S., Arai,

M., Kadowaki, K., “Neutron Powder Diffraction Study on Mg11B2 Synthesized by

Different Procedures”, J. Phys. Soc. Jap., 71(10), 2471-2476 (2002) (Experimental, Crys.

Structure, 49)

[2002Lee] Lee, S., Yamamoto, A., Mori, H., Eltsev, Yu., Masui, T., Tajima, S., “Single Crystals of

MgB2 Superconductor Grown under High-Pressure in Mg-B-N System”, Physica C,

378-381, 33-37 (2002) (Crys. Structure, Experimental, 10)

[2002Mir] Mironov, A., Kazakov, S., Jun, J., Karpinski, J., “MgB9N, a New Magnesium

Nitrido-Boride”, Acta Crystallogr., C58, i95-i97 (2002 (Experimental, Crys. Structure, 15)

[2003Kar] Kappinski, J., Kazakov, S. M., Jun, J., Angst, M., Puzniak, R., Wisniewski, A., Bordet, P.,

“Single Crystal Growth of MgB2 and Thermodynamics of Mg-B-N System at High

Pressure”, Physica C, 385, 42-48 (2003) (Experimental, Equi. Diagr., 17)

[2003Lee] Lee, S., “Crystal Growth of MgB2”, Physica C, 385(1-2), 31-41 (2003) (Crys. Structure,

Electr. Prop., Experimental, Magn. Prop., Phys. Prop., Review, Superconduct.,

Thermodyn., 56)

[2003Rec] Record, M.C., Tedenac, J.C., “B-N (Boron-Nitrogen)”, MSIT Binary Evaluation Program,

in MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services, GmbH,

Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram,

Assessment, 50)

107


MSIT®

B–Mg–N


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments /References

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

[Mas2]

( B)

< 2092

hR333

R3m

B

a = 1093.30

c = 2382.52

[1993Wer]

( B)

< 1000

hR36

R3m

B

a = 490.8

c = 1256.7

[V-C2]

low temperature - irreversible transition

to B

MgB2

1550 (B.P.)

hP3

P6/mmm

AlB2

a = 308.542

c = 351.97

a = 308.294

c = 351.348

a = 294

c = 322

a = 308.492

c = 352.007

[Mas2]

at 25°C [2002Oik];

range measured from 8 to 305 K

at 32 K [2002Oik]

at 39 GPa [2001Bor], read from graph

covering range up to 39 GPa.

at 5 GPa, 1600°C, 30 min [2002Lee]

MgB4

1775 (B.P.)

oP20

Pnma

MgB4

a = 546.4

b = 442.8

c = 747.2

[V-C2]

MgB7

2150 (B.P.)

oI64

Imma

MgB7

a = 597.0

b = 1048.0

c = 812.5

[V-C2]

lP - Mg3N2

< 897 at 5.5 GPa

cI80

Ia3

anti-Mn2O3

a = 996.44

a = 996.0

at 27°C, [2001Pas], range measured

from 11 to 305 K

[1993Lor]

hP-Mg3N2

897 - 1570; 5.5 GPa

Tetragonal a = 911.8

c = 669.4

at 1087°C, 5.5 GPa [1993Lor]

BNhex

“graphitic”

hP4

P63/mmc

BN

a = 250.428

c = 665.62

[1995Kur], 25°C

BNcubic

“sphalerite”

cF8

F43m

ZnS

a = 361.53 [1995Kur], 25°C

BN

“wurtzite”

hP4

P63mc

ZnS

a = 255.05

c = 421.0

[1995Kur], extrapolation to 1175°C

BN

“rhombohedral”

hR6 a = 250.42

c = 999

[1995Kur], 25°C

BN

“compressed hBN”

mC*

C2/c or Cc?

mBN

a = 433

b = 250

c = 310 to 330

= 92 to 95°

at 7.7 GPa, 1800-2150°C

[1996Hor]

108


MSIT®

B–Mg–N

B25N tP62

P42m

B25N

a = 863.4

c = 512.8

[V-C2]

* 1, lP - Mg3BN3

< 1200

at RT stable up to

5.8 GPa [1993Nak1]

hP14

P63/mmc

I-Mg3BN3

a = 354.453

c = 1603.536

a = 303.7

c = 1601.4

a = 354.2

c = 1601

[1991Hir]; see also [1993Sha];

[1992Lor], stable up to 4.4 GPa at RT

prepared at 1200°C [1994Zhu],

HV = 7.1 GPa; exp.= 2.39 Mg m-3

* 1, hP - Mg3BN3

< 1489 at 2.5 GPa

< 1685 at 6.5 GPa

oP7

Pmmm

II-Mg3BN3

a = 309.3

b = 313.4

c = 770.0

[1993Hir],

earlier “Mg3B2N4” [1979End1,

1979End2], earlier “Mg3B2N4”

[1989Hoh],

earlier “ Mg3B2N4” [1981Ely],

assumed as tetragonal by [1994Zhu]

(a = 310.7, c = 770 pm)

exp = 2.75 Mg m-3 [1994Zhu]

HV = 9.9 GPa [1994Zhu]

* 2, Mg6BN5

> 1150 at 4.5 GPa

< 1400 at 5.3 GPa

[1993Nak1]

hexagonal a = 539.7

c = 1058.5

[1999Kul1], prepared at 1600°C,

2.0 GPa; exp = 2.88 Mg m-3;

“X phase” of [1986Yam, 1993Nak1];

“hP-Mg3BN3” of [1989Hoh];

“ Mg3B2N4” of [1981Ely], [1994Zhu]

* 3, Mg3B2N4

hexagonal a = 1339.45

c = 595.17

at 5.5 GPa, >1277°C [1995Lor1,

1995Lor2, 1997Lor]

[1999Kul2]; exp = 2.67 Mg m-3

* 4, MgB9N hR66

R3m

MgB9N

a = 549.60

c = 2008.73

[2002Mir], prepared at 1600°C, 0.1 GPa.

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments /References

109


MSIT®

B–Mg–N

Temperature, °C

Pre

ssu

re,

GP

a

0 200 400 600 800 1000 1200 1400 1600

1

2

3

4

5

6

7

lP-Mg BN3 3

hP-Mg BN3 3

L

20

40

60

80

20 40 60 80

20

40

60

80

Mg B

N Data / Grid: at.%

Axes: at.%

BN

Mg3N

2

τ2

MgB2 MgB

4 MgB7

τ4

τ3

τ1

(Mg)

(αB)

Fig. 2: B-Mg-N.

Phase relations at

about 400°C. The

square denotes the

position of the

hP-Mg3B2N4 phase.

A half-filled circle -

position of the

MgB9N phase, filled

square - position of

the Mg6BN5 phase

Fig. 1: B-Mg-N.

p-T diagram for

Mg3BN3 with

experimental points

from [1993Nak1].

Solid circles are

lP-Mg3BN3, open

circles are

hP-Mg3BN3. The

dash-dotted curve is

from [1994Zhu]

110


MSIT®

B–Mg–N

Tem

pera

ture

,°C

Mg

N

B

Mg BN3 3Mg N3 2

L

1295°C

737°C

1489°C

BN

Mg BN + BN3 3

Mg BN + L3 3

Mg + Mg BN3 3

Mg BN + L3 3

Mg + L

800°C

BN + L

N, at.%

B, at.%

Mg, at.%

τ2

τ2

τ3

Mg

N

B

1227°C

~

~

1417°C

τ1= Mg BN3 3

τ3 = Mg B N3 2 4

BN

Mg N3 2τ1

τ3 + L

τ1+ L

BN + L

Tem

pera

ture

,°C

(BN + L)

ττ

1+

3 τ3 + BN

τ1 + BN

N, at.%

B,at.%

Mg, at.%

τ3

τ2 = Mg BN6 5

Fig. 3: B-Mg-N.

Temperature -

composition diagram

at 2.5 GPa for the

sections Mg-Mg3BN3

and Mg3BN3-BN,

after [1979End2]

modified for correct

composition of

compound

“Mg3B2N4” (now

Mg3BN3)

Fig. 4: B-Mg-N.

Temperature -

composition diagram

at 5.5 GPa for the

section Mg3BN3-BN

after [1997Lor]. The

range of existence for

the phase Mg6BN5

[1993Nak1,

1999Kul2] is added

111


MSIT®

B–Mg–N

1: reaction Mg N +hBN3 2

2: eutectic Mg BN3 3 3+hBN3: reaction Mg B N +hBN3 2 4

1

2

3

Pressure, GPa

Tem

pera

ture

,°C

2

727

8 10

927

1127

1327

1527

4 6

Fig. 5: B-Mg-N.

p-T reaction diagram

for the BN-Mg3N2

system after

[1994Gla]

1: [1979End1] (2%); 2: [1993Boc] (0.2%); 3: [1993Boc] (4%); 4: [1979End1] (8%);

5: [1966Kud]; 6: [1991Bin1, 1991Bin2] (9%); 7 and 8: various equilibrium lines

hBN-cBN (see [1993Boc]). The newly established line hBN-cBN after [2000Wil] is

outside the shown window: 2 GPa -2000°C, 4.5 GPa - 2750°C and 7 GPa - 3500°C

Temperature, °C

Pre

ssure

,G

Pa

1400

5

1200 1600 1800 2000 2200 2400 2600

6

71 2 3 4 5

6

7

8

Fig. 6: B-Mg-N.

p-T diagram for

growth of cBN in the

Mg-BN system and

role of oxygen. cBN

exists in the region

above each curve

(essentially after

[1993Boc]. Oxygen

contents of starting

BN materials is given

in parentheses

112


MSIT®

B–Mg–N

Temperature, °C

Pre

ssure

,G

Pa

800

5

600 1000 1200 1400 1600 1800 2000

6

71

23

4

5

6

7

4

Fig. 7: B-Mg-N.

p-T diagram for

growth of cBN in the

Mg3N2 - BN system

(after [1993Boc]).

cBN exists in the

region above each

curve

1: [1981Ely]; 2: [1981Sat]; 3: [1993Boc]; 4: [1972Dev]; 5: [1986Yam].

Curves 6, 7 are equilibrium lines for the conversion hBN-cBN (see [1993Boc])

113


MSIT®

B–N–Ti

Boron – Nitrogen – Titanium

Vasyl´ Tomashik

Literature Data

Critical assessments of the B-N-Ti ternary system have been published by [1996Dus], [1991Dus] and

[1992Rog], which cover the literature data up to 1991. Subsequently, this system has been investigated

experimentally using several techniques and for different temperature and composition ranges, and also has

been calculated thermodynamically.

Phase relationships in the B-N-Ti system appeared for the first time in the works of [1955Sam] and

[1955Bre]. The first experimental investigations of this ternary system were included in the review by

[1972Med]. According to [1955Sam] the solubility of TiB2 in TiN1-x reaches 8 mol% and the solubility of

TiN1-x in TiB2 is negligible. The crystal structure and lattice parameters of these solid solutions were

determined by [1971Aiv1, 1971Aiv2, 1971Aiv3, 1975Aly]. Such solid solutions can be obtained by

crystallization from the gas phase on the reduction of TiCl4-BCl3 mixtures under a nitrogen atmosphere.

There are some discrepancies between the experimental investigations of boron solubility in TiN1-x.

According to the data of [1971Aiv2, 1971Aiv3, 1973Tro] such solubility is too high and reaches 23.3 at.%.

However, [1975Aly] indicates that B solubility in TiN1-x at 1500°C is less then 1 at.%. The last value was

confirmed by further experimental and theoretical investigations of the B-N-Ti ternary system. As can be

seen from a comparison of the unit cell dimensions of binary and ternary phases, there is no significant solid

solubility of Ti in BN up to 1500°C, and mutual solubilities of the titanium borides, the titanium nitrides

and BN up to 1500°C are rather restricted [1996Dus].

[1981Chu1] and [1981Chu2] reported the existence of a pseudobinary section of the eutectic type for the

system TiB2-TiN revealing small mutual solid solubilities at the nitrogen-rich phase boundary TiN0.96,

whereas the solubility of TiB2 in TiN0.58 was said to increase up to ~12 mol% at 2300°C. The interaction

between titanium and BN results in a mixture of three phases TiB, TiB2 and TiN [1973Sam]. The solid state

reaction between Ti and BN powder begins at 1200°C (at 840-1100°C depending on the initial physical state

of the mixtures [1982Evt, 2001Gor]) and results in the formation of solid solutions of boron and nitrogen

in titanium and titanium borides and nitrides [1982Evt, 1984Bor, 2000Far]. The major part of the reaction

zone comprises the ( Ti) solid solution with grains containing fine Ti2N/( Ti) precipitates in a lamellar

structure formed during cooling from annealing temperature (1000-1200°C) [2000Far]. The phase

sequences at the interfaces are in good agreement with ternary B-N-Ti equilibrium diagram. The sequence

of layers in the coatings could be described as BN-TiB2-TiB-TiN1-x-( Ti)-Ti(pure) for the layers separated

by flat interfaces [2000Far]. The reaction between BN particles and the surrounding dense titanium matrix

at 1000°C yield a slightly different BN-TiB2-[TiB+TiN1-x]- Ti(N)-Ti(pure) phase sequence.

The combination of BN and TiN may decompose at high temperature and low partial pressure of nitrogen

according the following reaction: 2BN+TiN TiB2+3/2N2 [2001Rog].

A thermodynamic analysis of the reaction of Ti with BN suggests that self-propagating high-temperature

synthesis can be realized starting from ~9 mol% BN [2001Gor]. In this case, Ti based materials can be

obtained with different contents of TiB+TiN or TiN+TiB+TiB2, depending on the ratio of the starting

materials.

By studying the thermodynamics of the reaction 2BN+TiN TiB2+3/2N2 under 0.5 105 Pa of nitrogen and

related experimental investigations, [1955Bre] suggested that a mixture of TiN1-x+BN was stable up to

~1600°C. There was no evidence of solubility between the two in this pseudobinary system. Above

~1600°C, BN and TiN1-x will react to produce TiB2 and N2. The reported isothermal section of the B-N-Ti

phase diagram reveals a very limited solubility of B in TiN1-x and no solubility of nitrogen in TiB or TiB2

at 1400°C. The section is characterized by a dominating three phase field of TiB2+TiN1-x+BN [1961Now].

General agreement exists on the absence of ternary compounds in the B-N-Ti system.

By using a range of techniques which give direct information on crystalline structure, bonding types and

local atomic coordination and symmetry, [1997Mol] has demonstrated that there is a composition in the

114


MSIT®

B–N–Ti

B-N-Ti system for which PVD-synthesized thin films, deposited under certain conditions, do not exhibit the

structurally ordered phase mixture predicted by the accepted phase diagram for the bulk material. A large

fraction of the Ti atoms are situated in relatively disordered sites of lower symmetry than expected from the

crystalline material. The authors attribute the non-formation of the expected Ti containing phases for the

composition TiB1.7N1.8 to the combined effect of the high quenching rate associated with the production of

such films and the relatively low concentration of titanium. Taking into account all of experimental

information, they believe that the TiB1.7N1.8 material probably consists of clusters of atoms with varying

compositions and varying local symmetries, representing neither a fully nanocrystalline nor a

homogeneously amorphous state.

Phase equilibria in the B-N-Ti system have been investigated by [1987Smi] at 1500°C under high vacuum,

105 Pa of Ar and under 10 Pa N2, respectively. Phase relations at 1090°C have been determined by

[1991Dus]. Both of these isothermal sections were included in the reviews of [1996Dus, 1992Rog,

1994McH].

Binary Systems

The B-N system is accepted from [2003Rec]. Only one intermediate phase BN exists in this system. Boron

nitride has four crystalline structural modifications: cubic (cBN), wurtzite (wBN), hexagonal (hBN) and

rhombohedral (rBN). In addition, there are two other ordered BN phases: EBN, obtained by explosion (E)

of a mixture of hBN and aBN, compressed hBN attributable to a monoclinic lattice distortion of hBN and

two disordered BN phases: turbostratic BN (tBN) and amorphous BN (aBN).

The B-Ti system is accepted from [Mas2]. The mutual solid solubility of Ti and B is small (not higher than

1 at.%). TiB2 melts congruently whereas the TiB and Ti3B4 solids are incongruently melting.

The N-Ti system is taken from [Mas2]. The solubility of nitrogen in both ( Ti) and ( Ti) is significant. A

congruently melting TiN1-x compound having a wide region of homogeneity and an incongruently melting

Ti2N compound exist in this binary system. However, the phase diagram of [Mas2] is amended following

[1996Dus] and [1992Rog] who suggest two new phases, Ti3N2-x and Ti4N3-x form in the N-Ti system.

Solid Phases

No ternary compounds have been found in the B-N-Ti system. All unary and binary phases are listed in

Table 1.

Pseudobinary Sections

It is possible that the section TiB2-TiN1-x for x = 0.42 given by [1981Chu1, 1981Chu2] may be

pseudobinary (see below), but there is no experimental evidence indicating that the tie lines lie in the plane

of the section.

Isothermal Sections

The section of the B-N-Ti system at 200°C can be divided into three compatibility triangles (B-BN-Ti,

BN-Ti-TiN and BN-TiN-N). There is no solid solubility of the third component in any of the binary

compounds [1994McH]. Phase equilibria in the B-N-Ti system at 1090 and 1500°C (the former was

constructed by [1991Dus] and the latter was obtained under 100 kPa argon by [1961Now] and confirmed

by [1987Smi]) have been established from X-ray powder diffraction analysis and are given in Figs. 1 and 2,

respectively. These equilibria are characterized by the absence of ternary compounds and by the

incompatibility of titanium metal and hexagonal BN indicated by the presence of a stable tie line

TiB2-TiN1-x at temperatures below 1500°C.

A comparison of the unit cell dimensions shows no significant solubility of Ti in BN up to 1500°C. The

mutual solubilities of the titanium borides, the titanium nitrides and BN up to 1500°C are rather restricted

[1987Smi, 1991Dus, 1992Rog].

115


MSIT®

B–N–Ti


Three isopleths TiB2-TiN1-x for x = 0.04, 0.27 and 0.42 are presented in Fig. 3 [1981Chu1, 1981Chu2]

revealing the eutectic nature of a possible pseudobinary section between TiB2 and TiN1-x, and the increasing

solubility of TiB2 in TiN1-x with increasing x and temperature. These isopleths are included in the reviews

[1996Dus, 1991Dus, 1992Rog, 1994McH].


It was determined that maximum flexure strength and wear resistance at minimum constant of friction are

exhibited by eutectic alloys of the TiB2-TiN1-x sections [1983Tka], which were obtained according to the

procedure of [1981Chu1]. The experimental results indicate that TiN1-x is a good diffusion barrier for

boron; it allows limited diffusion of B in silicon at temperatures of up to 1000°C [1984Tin].

According to the data of [2000Bel], the addition of TiB2 as a reinforcing phase to TiN1-x based composites

improved both their hardness and strength in comparison to pure TiN1-x ceramic. The addition of TiB2 to

TiN1-x powder allowed high density to be achieved at lower temperatures and to limit grain growth.

The metallic nature of the coating formed on the surface of BN annealed in a loose Ti powder can provide

the surface metallization necessary for the successful joining of BN ceramic parts to metals and alloys

[2000Far].

Miscellaneous

All experimental data concerning the mutual solubility of titanium borides and nitrides can be summarized

in the tentative diagram presented in Fig. 4 [1989Bec], which was included in the review [1994McH] (in

the presented diagram the influence of N2 pressure has not been expressed). It can be seen that only a small

amount of boron can be incorporated in TiN1-x owing to its highly defective nitrogen sublattice. However,

in superstoichiometric Ti(N,B)1+x a considerable amount of B can be introduced, leading to interstitial solid

solutions, as has been suggested earlier by [1971Aiv2, 1971Aiv3, 1973Tro]. Slightly substoichiometric

TiN1-x is in equilibrium with nearly pure TiB2, therefore no reactions occur in compacts of TiN1-x and TiB2.

Under combustion conditions, fine-grained materials of high density (93 % to 95 %) can be obtained from

the BN-Ti mixtures [2001Gor]. When using a mixture 2Ti+BN, TiB and TiN1-x are distributed uniformly

in the Ti matrix (2Ti+BN TiB+TiN), and in the case of 3Ti+2BN mixtures, alloys of TiB+TiN with traces

of TiB2 are obtained (3Ti+2BN 2TiN+TiB2).

References

[1955Bre] Brewer, L., Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”,

J. Electrochem. Soc., 102(7), 399-406 (1955) (Experimental, Equi. Diagram, 19)

[1955Sam] Samsonov, G.V., Petrash, E.V., “Some Physico-Chemical Properties of Titanium Boride an

Nitride Alloys” (in Russian), Metalloved. Term. Obra. Metallov, (4), 19-24 (1955)

(Experimental, Equi. Diagram, 10)

[1961Now] Nowotny, H., Benesovsky, E., Brukl, C., Schob, O., “The Ternary Systems Ti-B-C and

Ti-B-N” (in German), Monatsh. Chem., 92(2), 403-414 (1961) (Experimental, Equi.

Diagram, #, *, 24)

[1971Aiv1] Aivazov, M.I., Domashnev, I.A., “Electrophysical Properties of Titanium Diboride and

Alloys in the System Ti-B-N”, Inorg. Mater., 7(10), 1551-1553 (1971), translated from Izv.

Akad. Nauk SSSR., Neorg. Mater., 7(10), 1735-1738 (1971) (Experimental, Equi.

Diagram, 7)

[1971Aiv2] Aivazov, M.I., Domashnev, I.A., Kireeva, I.M., “Electrical Properties of TiN0.96,

TiB0.43N0.78 and TiSi0.51N0.42” (in Russian), Izv. Akad. Nauk SSSR., Neorg. Mater., 7(10),


[1971Aiv3] Aivazov, M.I., Gurov, S.V., Domashnev, I.A., Kireeva, I.M., “Investigation of Magnetic

Properties of the Phases with Changeable Compositions as Titanium Nitride, Titanium

116


MSIT®

B–N–Ti

Diboride and the Alloys in the System Ti-B-N” (in Russian), Izv. Akad. Nauk SSSR., Neorg.

Mater., 7(7), 1176-1179 (1971) (Experimental, Equi. Diagram, Crys. Structure, 8)

[1972Med] Medvedeva, O.A., “System Metal-Boron-Nitrogen”, Sov. Powder Metall. Met. Ceram., 2,

113-118 (1972), translated from Poroshk. Metall., (2), 38-45 (1972) (Review, Equi.

Diagram, 28)

[1973Sam] Samsonov, G.V., Burykina, A.L., Medvedeva, O.A., Kosteruk, V.P., “The Interaction of

Boronitride with Transition Metals, their Borides and Nitrides”, Sov. Powder Metall. Met.

Ceram., 11, 903-908 (1973), translated from Poroshk. Metall., (11), 50-57 (1973)


[1973Tro] Troitsky, V.N., Grebtsov, B.M., Aivazov, M.I., “Obtaining of Titanium Boronitride

Powders in the Plasma SHF (Super High Frequency) Discharge” (in Russian), Poroshk.

Metall., (11), 6-9 (1973) (Experimental, Equi. Diagram, 6)

[1975Aly] Alyamovsky, S.I., Zainulin, Yu.G., Shveikin, G.P., Geld, P.V, Bausova, N.V., “Lattice

Defects in Cubic (NaCl Type) Zirconium and Titanium Boronitrides”, Inorg. Mater., 11(1),

148-149 (1975), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 11(1), 175- 176

(1975) (Experimental, Equi. Diagram, 12)

[1981Chu1] Chupov, V.D., Unrod, V.I., Ordanyan, S.S., “Reactions in the TiN-TiB2 System”, Soviet

Powder Metall. Met. Ceram., 1, 49-52 (1981), translated from Poroshk. Metall., (1), 62-66,

(1981) (Experimental, Equi. Diagram, #, *, 9)

[1981Chu2] Chupov, V.D., Ordanyan, S.S., Kozlovskii, L.V., “Interaction in the System TiNx-TiB2”,

Inorg. Mat., 17(9), 1195-1198 (1981), translated from Izv. Akad. Nauk SSSR., Neorg.

Mater., 17(9), 1618-1622 (1981) (Experimental, Equi. Diagram, #, *, 11)

[1982Evt] Evtushok, T.M., Zhunkovsky, G.L., “Contact Interaction of Titanium with Boron Nitride”

(in Russian), Zashch. Pokrytiya. Met., (16), 93-96 (1982) (Experimental, Equi. Diagram, 4)

[1983Tka] Tkachenko, Yu.G., Ordanyan, S.S., Yurchenko, D.Z., Yulyugin, V.K., Chupov, D.V.,

“High-Temperature Rubbing of the Alloys in the System TiNx-TiB2”, Sov. Powder Met.

Met. Ceram., 2, 137-141 (1983), translated from Poroshk. Metall., (2), 70-76 (1983)

(Experimental, Mechan. Prop., 6)

[1984Bor] Borisova, A.L., Borisov, Yu.S., Shvedova, L.K., Martsenyuk, N.S., “Interaction in Powder

Compositions Ti-BN”, Sov. Powder Metall. Met. Ceram., 4, 273-276 (1984), translated

from Poroshk. Metall., (4), 18-22 (1984) (Experimental, Equi. Diagram, 7)

[1984Tin] Ting, C.Y., “TiN as a High Temperature Diffusion Barrier for Arsenic and Boron”, Thin

Solid Films, 119(1), 11-21 (1984) (Experimental, Phys. Prop., 9)

[1987Smi] Smid, I., “Structural and Metallurgical Investigations in Boride and Boronitride Systems”

(in German), Thesis, University Vienna, 1-93 (1987) (Experimental, Equi. Diagram, #, *,

46) as quoted by [1996Dus, 1991Dus, 1992Rog]

[1989Bec] Becht, J.G.M., van der Put, P.J., Schoonman, J., “Chemical Vapor Deposition in the System

Ti-N-B: TiN as a Diffusion Barrier for Boron”, Europ. J. Solid State Inorg. Chem., 26(4),

401-412 (1989) (Review, Equi. Diagram, 25)

[1991Dus] Duschanek, H., Rogl, P., “The Ternary System Titanium-Boron-Nitrogen”, Leuven

Proceedings, COST 507, New Light Alloys, Effenberg, G. (Ed.), Part A, Belgium, A2, 1-9,

(1991) (Assessment, Experimental, Equi. Diagram, #, *, 27)

[1992Rog] Rogl, P., Schuster, J.C., “Ti-B-N (Titanium-Boron-Nitrogen)”, Phase Diagrams of Ternary

Boron Nitride and Silicon Nitride Systems. Monogr.Ser. of Alloy Phase Diagr., 103-106

(1992) (Review, Equi. Diagram, Crys. Structure, Thermodyn., #, *, 19)


of Carbon-Doped -Rhombohedral Boron”, Phys. Status Solidi (B), B179, 489-511 (1993)


[1994McH] McHale, A.E., “VIII. Boron+Nitrogen+Metal; B-N-Ti”, Phase Equilibria Diagrams, Phase

Diagrams for Ceramists, 10, 238-240 (1994) (Review, Equi. Diagram, 7)

[1996Dus] Duschanek, H., Rogl, P., “Boron-Nitrogen-Titanium”, MSIT Ternary Evaluation Program,

in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services

117


MSIT®

B–N–Ti


Assessment, 25)

[1997Mol] Mollart, T.P., Gibson, P.N., Baker, M.A., “An EXAFS and XRD Study of the Structure of

Nanocrystalline Ti-B-N Thin Films”, J. Phys. D: Appl. Phys., 30, 1827-1832 (1997)


[2000Bel] Bellosi, A., Monteverde, F., “Microstructure and Properties of Titanium Nitride and

Titanium Diboride-Based Composites”, Key Eng. Mater., 175-176, 139-148 (2000)

(Experimental, Mechan. Prop., Phys. Prop., 57)

[2000Far] Faran, E., Gotman, I., Gutmanas, Y., “Experimental Study of the Reaction Zone at Boron

Nitride Ceramic - Ti Metal Interface”, Mater. Sci. Eng. A, A288, 66-74 (2000)


[2001Gor] Gordienko, S.P., Evtushok, T.M., “Reaction of Titanium with Boron Nitride under

Self-Propagating High-Temperature Synthesis Conditions”, Powder Metall. Met. Ceram.,

40(1-2), 58-60 (2001), translated from Poroshk. Metall., (1-2), 76-79 (2001) (Calculation,

Thermodyn., 3)

[2001Rog] Rogl, P., “Materials Science of Ternary Metal Boron Nitrides”, Int. J. Inorg. Mater., 3,


[2003Rec] Record, M.Ch., Tedenac, J.-C., “B-N (Boron-Nitrogen)”, MSIT Evaluation Program, in


GmbH, Stuttgart, to be published, (2002) (Review, Equi. Diagram, 50)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( B)

< 2092

hR333

R3m

B

a = 1093.30

c = 2382.52

pure B [Mas2, 1993Wer]

( N)

< -237.54

cP8

Pa3

N

a = 566.1 [Mas2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 [Mas2];

dissolves 23 at.% N at 1050C [Mas2]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

at 25°C [Mas2];

dissolves 6.2 at.% N at 2020°C [Mas2]

hBN

< 2397

hP4

P63/mmc

BN

a = 250.4

c = 666.1

[2003Rec]

cBN cF8

F43m

ZnS

a = 361.53 0.04 [2003Rec]

wBN hP4

P63/mc

ZnS

a = 255.0 0.5

c = 423 1

[2003Rec]

rBN hR6 a = 250.4

c = 999.1

[2003Rec]

118


MSIT®

B–N–Ti

Compressed hBN mC4

C2/c or Cc

a = 433

b = 250

c = 310 to 330

= 92-95°

[2003Rec]

TiB

< 2190

oP8

Pnma

FeB

a = 612 1

b = 306 1

c = 456 1

[V-C2]

Ti3B4

< 2200

oI14

Immm

Ta3B4

a = 325.9

b = 1373

c = 304.2

[V-C2]

TiB2

< 3225

TiB2.21N0.23

1430

TiB2.07N0.29

1530

TiB1.98N0.33

1630

hP3

P6/mmm

AlB2

a = 303.8

c = 323.9

a = 304

c = 323.1

a = 303.8

c = 322.5

a = 303.7

c = 322.3

[V-C2]

[1971Aiv1, 1971Aiv2, 1971Aiv3]

[1971Aiv1, 1971Aiv2, 1971Aiv3]

[1971Aiv1, 1971Aiv2, 1971Aiv3]

Ti2N

< 1100

tP6

P42/mnm

TiO2

a = 494.52

c = 303.42

at 33 to 34 at.% N

[V-C2]

TiN1-x

< 3290

TiB0.425N0.78

1230

TiB0.54N0.77

1230

TiB0.005N0.62

1500

TiB0.01N0.73

1500

TiB0.01N0.77

1500

TiB0.03N0.76

1500

TiB0.02N0.82

1500

TiB0.05N0.76

1500

cF8

Fm3m

NaCl

a = 423.9 0.1

a = 425

a = 423

a = 422.93

a = 423.44

a = 423.43

a = 423.76

a = 423.60

a = 423.82

[V-C2]

[1971Aiv2, 1971Aiv3]

[1971Aiv2]

[1975Aly]

[1975Aly]

[1975Aly]

[1975Aly]

[1975Aly]

[1975Aly]

Ti3N2-x

1103 - 1066

hR2

? VTa2C2

a = 297.95

c = 2896.5

at 29 at.% N

[1996Dus, 1992Rog]

Ti4N3-x

1291 - 1078

hR2

? V4C3

a = 298.09

c = 2166.42

at 31.5 at.% N

[1996Dus, 1992Rog]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

119


MSIT®

B–N–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti B

N Data / Grid: at.%

Axes: at.%

BN

TiB2Ti

3B

4TiB

(βTi)

(αTi)

Ti3N

2-x

Ti4N

3-x

δ, TiN1-x

(αTi)+(βTi)+TiB

Ti3B

4+TiN

1-x+TiB

2

Ti4N

3-x+Ti

3N

2-x+TiB

(βB)

TiN+BN+N2

(αTi)+Ti3N

2-x+TiB TiB

2+BN+(βB)

TiN1-x

+Ti4N

3-x+TiB

TiN1-x

+BN+TiB2

Fig. 1: B-N-Ti.


1090°C under 105 Pa

of argon (in the

absence of external

nitrogen)

20

40

60

80

20 40 60 80

20

40

60

80

Ti B

N Data / Grid: at.%

Axes: at.%

(βTi)

(αTi)

TiN1-x

BN

TiB2

Ti3B

4TiB

TiN1-x

+Ti3B

4+TiB

TiN1-x

+Ti3B

4+TiB

2

TiB2+BN+(βB)

TiN1-x

+(αTi)+TiB

TiN+BN+TiB2

TiN+BN+N2

(αTi)+(βTi)+TiB

(βB)

Fig. 2: B-N-Ti.


1500°C under 105 Pa

of argon (in the

absence of external

nitrogen)

120


MSIT®

B–N–Ti

Ti 33.33

B 66.67

N 0.00

Ti 51.00

B 0.00

N 49.00N, at.%

Te

mp

era

ture

, °C

TiN0.58+L2335±40°C

2447±40°C

2600±45°C

2935±70°C

3057±80°C

TiN0.73+L

TiN0.96+L

L

L

L

20

40

60

80

20 40 60 80

20

40

60

80

Ti B

N Data / Grid: at.%

Axes: at.%

TiB TiB2

TiN

Fig. 3: B-N-Ti.

Comparison of the

concentration

sections TiB2-TiN1-x

at x = 0.04, 0.27 and

0.42 under 10 MPa of

N2, Ar

Fig. 4: B-N-Ti.

Tentative diagram

showing the

subsolidus phase

relationships

121


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C–N–Ti

Carbon – Nitrogen – Titanium

Volodymyr Ivanchenko

Literature Data

The C-N-Ti system is of great technical importance due to the formation of a solid solution between TiC

and TiN exhibiting high hardness, high corrosion resistance and high thermodynamic stability. The titanium

carbonitrides are very often used for bulk strengthening as well as surface treatment of different materials.

Conversely, carbon and nitrogen embrittle titanium, when dissolved. For this reason, experimental

investigation of phase equilibria in the C-N-Ti system began in the early 50s, during the first stage of

development of the titanium industry.

The phase relations between titanium monocarbide and mononitride were first studied by [1950Duw]. It

was shown that TiC and TiN formed a continuous series of solid solutions and the lattice parameter vs

composition curve was almost linear. [1953Sto] investigated phase equilibria in the Ti corner of the phase

diagram at 800, 900, 1000 and 1300°C for carbon contents of up to 7 at.% and nitrogen contents of up to

12 at.%, mainly by metallography. Phase equilibria were established in the Ti rich corner because the alloys

containing more than 3.6 at.% C and more than 9.3 at.% N were inhomogeneous. [1972Kli] synthesized

specimens of titanium carbonitride with different C/(C+N) ratios and roughly determined the homogeneity

region of Ti(CxN1-x)y. [1978Arb1] presented the isothermal section at 500°C which included the phase

equilibria involving the ordered phases in the C-Ti and N-Ti binary systems. The solid state phase equilibria

at 1150°C were investigated by analyzing arc melted and hot pressed alloys for nitrogen and carbon. X-ray

diffractometry, metallography and electron probe microanalysis were used. A tentative isothermal section

at 1150°C was presented by [1995Bin].

The crystal structures of carbonitrides were examined by [1974Mit, 1976Vil, 1978Arb2, 1994Aig,

1998Wok] using X-ray diffractometry and neutron diffraction [1976Kar, 1987Em, 1996Tas].

The heat content of Ti(CxN1-x) has been measured between 227 and 1227°C by [1982Tur, 1984Tur] using

solution calorimetry. The first Calphad assessment was proposed by [1984Tey].

[1971Bog, 1975Zhi] investigated physicochemical properties of titanium carbonitrides such as solubility in

different acids, high-temperature oxidation and microhardness. The elastic properties of Ti carbonitrides of

nonstoichiometric compositions were examined by [1976Iva]. Thermal conductivity, electrical

conductivity and thermal expansion of titanium carbonitrides were presented by [1978Iva]. Technically

relevant solid state properties of the hot-pressed carbonitrides Ti(CxN1-x)~1.00 and Ti(CxN1-x)0.82, such as

microhardness, electrical conductivities, heat conductivities and optical reflectance were measured as

functions of the x =C/(C+N) ratio by [1995Len]. The lattice parameters and thermal expansion coefficients

across the whole range of TiN-TiC compositions over the temperature range 25-1200°C were determined

on polycrystalline specimens by [1994Aig] and a polynomial fitting = f(x) was proposed. The same

properties were measured between 25 and 327°C on whiskers and needle like crystals by [1998Wok].

[1981Arb] studied the abrasive ability of titanium carbonitride. Synthesis of titanium carbonitride was

elaborated using a solar furnace [1999Fer] and chemical vapor deposition at a moderate temperatures

[2002Lar].

A method for calculating the nitrogen equilibrium partial pressure pN2 of stoichiometric and

nonstoichiometric carbonitrides having the NaCl type structure has been proposed by [1995Gus, 1996Gus,

1998Gus], and nitrogen partial pressure as a function of temperature and composition for TiNy

(0.45 y 1.0) and TiCxNy (0.50 x+y 1.0, 0 y 1.0) were calculated and compared with experimental

data from the literature.

The results of spectroscopic and theoretical investigations of substoichiometric titanium nitrides and

carbonitrides were presented by [1997Gue]. The phase stability of Ti(C,N) solid solutions was theoretically

investigated by [1999Jun].

Thermodynamic calculations for the C-N-Ti system based on different assessments of the N-Ti and C-Ti

binary systems have been performed by [1984Tey, 1993Oht, 1996Jon, 1999Dum, 2001Lee].

122


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C–N–Ti

[1983Sch] briefly discussed the topology of the C-N-Ti system based on the results of [1953Sto] and

[1978Arb1].

An assessment of the C-N-Ti system based on experimental studies and calculations of phase equilibria was

presented by [2000Ban].

Binary Systems

The assessed C-Ti phase diagram is presented in Fig. 1. It is the compilation of a thermodynamic

optimization by [1996Sei] and the results of a neutron diffraction study of the nonstoichiometric ordered

structures of Ti2C1+x that exist at lower temperatures, given by [1991Tas, 2002Tas]. [1996Sei] did not take

into account the ordering of carbon on the carbon + vacancy sublattice at compositions near Ti2C in the

abscence of thermodynamic data. [2002Tas] established, by a high resolution neutron diffraction study of

the titanium carbides TiC0.59 and TiC0.62, that there are two stable ordered structures; one with a trigonal

unit cell and a second with a cubic unit cell. On decreasing temperature, ,TiCx (x~0,6) undergoes a second

order tranformation (Fm3m) ’(Fd3m) at 790°C and a first order transformation ’(Fd3m) ’’(P3121)

at 770°C. The thermodynamic assessment of [1996Sei], reproduced by [1998Oka], does not take into

account the ordering of TiC observed below 790°C.

The assessed N-Ti phase diagram is taken from [1992Len] who modified the literature compilation of

[1987Wri] by including their own results based on diffusion couple investigations. This phase diagram was

presented by [1993Oka].

Solid Phases

The system shows a complete series of solid solutions between the isostructural binary phases TiC1-x and

TiN1-x. Table 1 summarizes the crystal structure data of the binary phases, including the ordered phases and

the complete solid solution Ti(CxN1-x)y.


An optimization of the TiN-TiC pseudobinary section was first performed by [1996Jon], then by

[1999Dum] for pN2 = 1 bar. New information on the C-Ti system was taken into account. The pseudobinary

section calculated by [1999Dum] is shown in Fig. 2.


Table 2 presents the only known invariant equilibrium, L+ + , which occurs at 1820°C [1996Jon].

Liquidus Surface

The liquidus surface shown in Fig. 3, is based mainly on the calculations of [1996Jon]. The diagram has

been modified to match the accepted the C-Ti and N-Ti binary systems.

Isothermal Sections

The isothermal section of the C-N-Ti phase diagram at 500°C is taken from [1978Arb1] and given in Fig. 4,

after modification to maintain consistency with the accepted binary systems. The tentative phase diagram

at 1150°C presented in Fig. 5 is taken from [1995Bin]. The isothermal section at 1820°C calculated by

[1996Jon] is presented in Fig. 6. It should be noted that the results of the calculations of [1996Jon] at 500,

1200 and 1700°C are similar to those of [1999Dum] despite some differences in the interaction parameters

used in the calculations. A good fit to the experimental information was obtained for the whole range of

compositions and temperatures, except for the composition of the carbonitride, where the experimental

scatter is very large.

123


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C–N–Ti

Thermodynamics

[1982Tur] measured the heat content of three alloys between 227 and 1227°C, and found a linear variation

from TiC to TiN. This suggests that the solid solution between TiC and TiN is nearly ideal. [1984Tur]

suggested that the alloys contained a mixture of TiC and TiN and repeated the measurements on four new

alloys between 227 and 1227°C. A distinct minimum in the enthalpy and heat capacity were observed for

the composition TiC0.66N0.28. The temperature and concentration dependence of enthalpy and entropy of

the TiCxN1-x carbonitrides were presented as a polynomial expansion in [1984Tur], as:

(HT-HT°)/J mol-1 = -16558 + 34.148x - 9.717xT + 45.584T + 3.971 10-3T2 + 8.981 105T-1 + 7.918 x2T 285

ST° = 104.98 log10 T + 7.942 10-3T + 4.490 105T-2 - 22.38x log10T + 18.22x2 log10T + 27.307x - 22.287 x2

- 236.97 J mol-1 K-1

The deviation of the calculated values of entropy from the experimental data was less than 1%.

[1995Len] measured the heat capacity between 50 and 1000°C using a sample with a composition given as

TiC0.6N0.4. The result agreed with [1984Tur] within experimental error.


The titanium carbonitrides can be dissolved in a mixture of HNO3+HF (1:1). The solubility of TiCxNy in

HCl is very low and is practically independent of composition. The stability of titanium carbonitrides in

HCl+HNO3 (3+1), H2SO4+HNO3 (1:1), HNO3 (concentrated) and HNO3 (1:1) decreases with increasing

carbon content [1971Bog].

Titanium carbonitrides with compositions close to TiC0.6N0.3 and TiC0.3N0.6 have higher oxidation

resistance above 660°C. But in the temperature interval 520 to 570°C, specimens of all compositions have

an anomalously high rate of oxidation. Specimens enriched in carbon first lose nitrogen, and specimens

depleted in carbon first lose carbon [1975Zhi].

The composition dependence of the elastic properties of titanium carbonitrides is not linear. On the

substitution of C by N in the TiCxNy solid solution the rigidity modulus at first decreases and reaches a

minimum at TiC0.61N0.31 before increasing to a maximum at y 0.6 [1976Iva]. Their room temperature

values lie between the limits E = 360-430 GPa, K = 190-210 GPa and G = 180-200 GPa.

The Vickers hardness at a load of 0.98 N was measured by [1995Len] as a function of the C/(C+N) ratio for

two carbonitride series, Ti(CxN1-x)1.00 and Ti(CxN1-x)0.82. For Ti(CxN1-x)y, the variation in the hardness

with composition for the stoichiometric and substoichiometric samples shows a positive deviation from

additive behavior. The measured values vary from 16.5-18 GPa (TiN1.0-0.82) to 25-30GPa (TiC0.82-1.0). The

data of [1971Bog] for titanium carbonitrides show a large deviation from the values presented by

[1995Len], except for the boundary compounds TiN and TiC, with a very high maximum of 38.2 GPa at a

composition of TiC0.57N0.39.

In accordance with [1995Len], the composition dependence of electrical conductivity at room temperature

may be fitted as

= 0.0389 - 0.0903 (C/(C+N)) + 0.1467 (C/(C+N))2 - 0.0868 (C/(C+N))3 (106 -1 cm-1).

The temperature conductivity and the heat conductivity increase with increasing nitrogen content. The data

points were fitted by the polynomial expansions:

a(T)=A+BT+CT2+DT-2 (102 cm2 s-1) for the temperature conductivity and as k(T)=A+BT+CT2+DT-2

(W m-1 K-1) for heat conductivity. The coefficients are presented in Tables 3 and 4 after [1995Len].

The average thermal expansion coefficient of polycrystalline TiCxN1-x in the range of 25-1200°C was

measured by [1994Aig] as av = (9.9-1.4 C/(C+N)) 10-6 K-1.

Miscellaneous

[1995Gus, 1996Gus, 1998Gus] proposed a method for calculation of the nitrogen partial pressure of

stoichiometric and nonstoichiometric nitrides and carbonitrides and presented the results of calculations for

TiCxNy (0.50 x+y 1.0, 0 y 1.0) in the temperature range 1327-1727°C. It was shown, that

nonstoichiometry of the carbonitride nonmetallic sublattice appreciably affects the value of nitrogen partial

pressure. A high nitrogen partial pressure will be observed over carbonitride with a composition closer to

124


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C–N–Ti

stoichiometric. For example, at 1727°C the value of pressure pN2 over carbonitrides TiC0.75N0.2 (x+y=0.95)

and TiC0.5N0.2 (x+y=0.7) are equal to 226 Pa and 0.068 Pa, respectively.

[1999Fer] studied the synthesis of titanium carbonitride from a compact mixture of pure metal powder and

carbon powder (graphite or amorphous carbon) using a solar furnace heating for 30 min at a measured

temperature around 1600°C. No evidence of graphitization of the amorphous carbon was observed. The

composition of the carbonitride depended on whether the carbon was graphitic or amorphous.

In the work of [2002Lar], Ti(C,N) coatings were produced by moderate temperature chemical vapor

deposition (MTSVD). MTCVD coated tools have higher transverse rupture strength values than those

coated via a CVD route owing to reduced decarburization of the cementite carbide substrates and lower

residual tensile stresses. Reduced decarburization and the absence of the reversible phase reaction lead to

improved edge strength.

References

[1950Duw] Duwez, P., Odell, F., “Phase Relationships in the Binary Systems of Nitrides and Carbides

of Zirconium, Columbium, Titanium and Vanadium”, J. Electrochem. Soc., 97(9), 290-297


[1953Sto] Stone, L., Margolin, H., “Titanium-Rich of the Ti-C-N, Ti-C-O, and Ti-N-O Phase

Diagrams”, J. Inst. Met., 5, 1498 (1953) (Experimental, Equi. Diagram, 10)

[1971Bog] Bogomolov, V.A., Shveikin, G.P., Alyamovskiy, S.I., Zainulin, Yu.G., Liubimov, V.D.,

“Physicochemical Properties of Titanium Oxinitrides and Carbonitrides” (in Russian), Izv.

Akad. Nauk SSSR, Neorg. Mater., 7(1), 67-72 (1971) (Experimental, 15)

[1972Kli] Klimashin, G.M., Avgustinnik, A.I., Smirnov, G.V., “About Carbonitride and Oxicarbide

Phases of Titanium and Zirconium” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater.,

8(5), 843-845 (1972) (Equi. Diagram, Experimental, 8)

[1974Mit] Mitrofanov, B.B., Zainulin, Yu.G., Alyamovskiy, S.I., Shveikin,G.P. “Region of

Homogeneity, Degree of Filling and Concentration Dependence of Lattice Parameters of

Cubic Titanium Carbonitride” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 10(4),

745-747 (1974) (Equi. Diagram, Crys. Structure, Experimental, 10)

[1975Zhi] Zhilyaev, V.A., Shveikin, G.P., Alyamovskiy, S.I., Liubimov, V.D., Mitrofanov, B.V.,

“Kinetic of High-Temperature Oxidation of Titanium Carbonitrides in the Air” (in

Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 11(2), 230-235 (1976) (Experimental, 3)

[1976Iva] Ivanov, N.A., Andreeva, L.P., Alyamovskiy, S.I., Mitrofanov, B.V., “The Elastic Properties

of Ti Carbonitrides of Nonstoichiometric Compositions” (in Russian), Izv. Akad. Nauk

SSSR, Neorg. Mater., 12(7), 1209-1211, (1976) (Experimental, Mechan. Prop.,

Experimental, 8)

[1976Kar] Karimov, I., Em, V.T., Petrunin, V.F., Latergaus, I.S., Polishchuk, V.S.,

“Neutron-Diffraction Study of Titanium Carbonitrides”, (in Russian) Izv. Akad. Nauk SSSR,

Neorg. Mater., 12, 1492-1494 (1976) (Crys. Structure, Experimental, 9)

[1976Vil] Vilk, Yu.N., Danisina, I.N., “Structural Parameters and X-Ray and Pycnometric Densities

of Ti Carbonitride” (in Russian), Poroshk. Met., (12), 42-48 (1976) (Crys. Structure,

Experimental, 25)

[1978Arb1] Arbuzov, M.P., Golub, S.Ya., Khaenko, B.V., “The Investigation of Phases in the

Ti-TiC-TiN System” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 14(8), 1442-1448

(1978) (Equi. Diagram, Crys. Structure, Experimental, #, 7)

[1978Arb2] Arbuzov, M.,P., Golub, S., Ya., Khaenko, B.,V., “A Distortion of the Crystal Lattice of the

Ordered Phase on the Base of Cubic Titanium Carbonitride”, Dop. Akad. Nauk Ukr. RSR,

Ser.A, Fiz-Mat. Tekh. Nauki, (in Ukrainian), (2), 181-183 (1978) (Crys. Structure,

Experimental, 5)

[1978Iva] Ivanov, N.A., Andreeva, L.P., Gel’d, P.V., “Heat Conductivity, Electrical Resistivity and

Thermal Expansion of Titanium Carbonitrides and Oxicarbides”, Poroshk. Metall., (8),

54-58 (1978) (Experimental, 15)

125


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C–N–Ti

[1981Arb] Arbuzov, M.P., Moshkovskiy, E.I., Liaschenko, A.B., “Abrasive Abilities of Titanium

Carbonitride”, (in Russian), Poroshk. Metall., (6), 78-81 (1981) (Experimental, 7)

[1982Tur] Turchanin, A.G., Babenko, S.A., Polischuk, V.S., “Enthalpy and Heat Capacity of Cubic

Titanium Carbonitrides in the Temperature Interval of 298-1500 K” (in Russian), Zh. Fiz.

Chim., 56(1), 41-44 (1982) (Thermodyn., Experimental, 16)

[1983Sch] Schouler, M.C., Ducarroir, M., Bernard, C., “Review of the Constitution and the Properties

of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System”, Rev. Int. Hautes Temp.

Refract., (in French), 20, 261-311 (1983) (Review, Equi. Diagram, 154)

[1984Tey] Teyssandier, F., Ducarroir, M., Bernard, C., “Thermodynamic Study of the

Titanium-Carbon-Nitrogen Phase Diagram at High Temperature”, Calphad, 8(3), 233-242

(1984) (Calculation, Equi. Diagram, Thermodyn., 25)

[1984Tur] Turchanin, A.G., Babenko, S.A., Bilyk, I.I., “Temperature and Composition Dependences

of Titanium Carbonitrides Thermodynamical Properties in Temperature Interval of

298-1500 K” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 20(9), 1511-1514 (1984)

(Thermodyn., Experimental, 7)

[1986Len1] Lengauer, W., “The Crystal Structure of a New Phase in the Titanium-Nitrogen System”,

J. Less-Common Met. , 120, 153-159 (1986) (Crys. Structure, Experimental, 24)

[1986Len2] Lengauer, W., “The Crystal Structure of -Ti3N2-x: An Additional New Phase in the Ti-N

System, J. Less-Common Met. , 125, 127-134 (1986) (Crys. Structure, Experimental, 19)

[1987Em] Em, V.T., Karimov, I., Latergaus, I.S., “Influence of Nitrogen on the Capability of

Order-Disorder Phase Transformation in Tix” (in Russian), Metallophyzika, 9(4), 113-114


[1987Len] Lengauer, W., Ettmayer, P., Some Aspects of the Formation of -Ti2N, Rev. Chim. Miner.,

24, 707-713 (1987) (Equi. Diagram, Crys. Structure, Experimental,13)

[1987Wri] Wriedt, H.,A., Murray, J.L., “The N-Ti (Nitrogen-Titanium) System”, Bull. Alloy Phase

Diagr., 8(4), 378-388 (1987), (Review, Equi. Diagram, Crys. Structure, Thermodyn., *, 56)

[1989Kha] Khaenko, B.V., Kukolí, V.V., “Real Structure of the Ordered Titanium Carbide” (in

Russian), Kristallographiya, 34(6), 1513-1517 (1989) (Crys. Structure, Experimental, 10)

[1991Tas] Tashmetov, M.Yu., Em, V.T., Kalanov, M.U., Shkiro, V.,M., “An Ordering Structure and

Phase Transformations in Titanium Carbide”, (in Russian), Metallofizika, 13(5), 100-106

(1991) (Crys. Structure, Experimental, *, 13)

[1992Len] Lengauer, W., “Properties of Bulk -TiN1-x Prepared by Nitrogen Diffusion into Titanium

Metal”, J. Alloys Compd., 186, 293-307 (1992) (Equi. Diagram, Crys Structure, #, 35)

[1993Oht] Ohtani, H., Hillert, M., “Calculation of V-C-N and Ti-C-N Phase Diagram”, Calphad,

17(1), 93-99 (1993) (Equi. Diagram, Thermodyn., Calculation, 15)

[1993Oka] Okamoto, H., “N-Ti (Nitrogen-Titanium)”, J. Phase Equilib., 14(4), 536 (1993)

(Assessment, Equi. Diagram, #, 8)

[1994Aig] Aigner, K., Lengauer, W., Rafaja, D., Ettmayer, P., “Lattice Parameters and Thermal

Expansion of Ti(xN1-x), Zr(CxN1-x), HfxN1-x and TiN1-x from 298 to 1473 K as Investigated

by High-Temperature X-Ray Diffraction”, J. Alloys Compd., 215, 121-126 (1994) (Crys.


[1995Bin] Binder, S., Lengauer, W., Ettmayer, P., Bauer, J., Debuigne, J., Bohn, M., “Phase Equilibria

in the Systems Ti-C-N, Zr-C-N and Hf-C-N”, J. Alloys Compd., 217, 128-136, (1995)

(Equi. Diagram, Crys. Structure, Experimental, #, 33)

[1995Gus] Gusev, A.I., “Nitrogen Pressure over Cubic Stoichiometric and Nonstoichiometric

Transition Metal Nitrides and Carbonitrides”, (in Russian), Dokl. Akad. Nauk SSSR, 340(6),

758-762 (1995) (Theory, Thermodyn., Calculation, 13)

[1995Len] Lengauer, W., Binder, S., Aigner, K., Ettmayer, P., Guillou, A., Debuigne, J., Groboth, G.,

“Solid State Properties of Group IVb Carbonitrides”, J. Alloys Compd., 217, 137-147 (1995)

(Review, Experimental, 46)

[1996Gus] Gusev, A.I., “The Influence of Composition, Nonstoichiometry, and Temperature on the

Partial Pressure of Nitrogen Over Metal Nitrides and Carbonitrides”, Russ. J. Phys. Chem.,

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C–N–Ti

70(4), 570-575 (1996), translated from Zh. Phys. Khim., 70(4), 616-621 (1996)

(Calculation, Thermodyn., 14)

[1996Jon] Jonsson, S., “Calculation of the Ti-C-N System”, Z. Metallkd., 87(9), 713-720 (1996) (Equi.

Diagram, Assessment, Thermodyn., #, 24)

[1996Tas] Tashmetov, M.Yu., Em, V.T., Mukhtarova, N.N., “Trigonal Ordered Structure in Titanium

Carbonitride”, Inorg. Mater., 32 (2), 151-152 (1996), translated from Neorg. Mater., 32(2),

171-172 (1996), (Crys. Structure, Experimental, 7)

[1996Sei] Seifert, H.J., Lukas, H.L., Petzow, G., “Thermodynamic Optimization of the Ti-C System”,

J. Phase Equilib., 17(1), 24-35 (1996) (Equi. Diagram, Thermodyn., Calculation, #, 58)

[1997Gue] Guemmaz, M., Moraitis, G., Mosser, A., Khan, M.A., Parlebas, J.C., “Band Structure of

Substoichiometric Titanium Nitrides and Carbonitrides: Spectroscopical and Theoretical

Investigations”, J. Phys.: Condens. Matter, 9, 8453-8463 (1997) (Crys. Structure,

Experimental, 27)

[1998Gus] Gusev, A.I., “Nitrogen Partial Pressure of Stoichiometric and Non-Stoichiometric

Titanium, Vanadium, and Niobium Nitrides and Carbonitrides”, Phys. Status Solidi B,

209(2), 267-286 (1998) (Thermodyn., 35)

[1998Oka] Okamoto, H., “C-Ti (Carbon-Titanium)”, J. Phase Equilib., 19(1) 89 (1998) (Assessment,

Equi. Diagram, 10)

[1998Wok] Wokulska, K., “Thermal Expansioin of Whiskers of Ti(C,N) Solid Solutions”, J. Alloys

Compd., 264, 223-227 (1998) (Crys. Structure, Experimental, 29)

[1999Dub] Dubrovinskaia, N.A., Dubrovinsky, L.S., Saxena, S.K., Ahuja, R., Johansson, B.,

“High-Pressure Study of Titanium Carbide”, J. Alloys Compd., 289, 24-27 (1999) (Crys.


[1999Dum] Dumitrescu, L.F.S., Hillert, M., Sundman, B., “A Reassessment of Ti-C-N based on a

Critical Review of Available Assessments of Ti-N and Ti-C”, Z. Metallkd., 90(7), 534-541

(1999) (Assessment, Thermodyn., #, 38)

[1999Fer] Fernandes, J.C., Amaral, P.M., Rosa, L.G., Martinez, D., Rodrigues, J., Shohoji, N., “X-Ray

Diffraction Characterisation of Carbide and Carbonitride of Ti and Zr Prepared Through

Reaction Between Metal Powders and Carbone Powders (Graphitic or Amorphous) in a

Solar Furnace”, Int. J. Refract. Mater., 17, 437-443 (1999) (Crys. Structure,

Experimental, 17)

[1999Jun] Jung, I.-J., Kang, S., Jhi, S.-H., Ihm, J., “A Study of the Formation of Ti(CN) Solid

Solutions”, Acta Mater., 47(11), 3241-3245 (1999) (Calculation, Thermodyn., 15)

[2000Ban] Bandyopadhyay, D., Sharma, R.C., Chakraborti, N., “The Ti-N-C System

(Titanium-Nitrogen-Carbon)”, J. Phase Equilib., 21(2), 192-194 (2000) (Assessment,

Thermodyn., Crys. Structure, #, 15)

[2001Lee] Lee, B.-J., “Thermodynamic Assessment of the Fe-Nb-Ti-C-N System”, Metall. Mater.

Trans. A, 32A(10), 2423-2439 (2001) (Assessment, Equi. Diagram, Thermodyn., 90)

[2002Lar] Larsson, A., Ruppi, S., “Microstructure and Properties of Ti(C,N) Coatings Produced by

Moderate Temperature Chemical Vapour Deposition”, Thin Solid Films, 402, 203-210

(2002) (Experimental, Mechan. Prop., 14)

[2002Tas] Tashmetov, M.Yu., Em, V.T., Lee, C.H., Shim, H.S., Choi, Y.N., Lee, J.S., “Neutron

Diffraction Study of the Ordered Structures of Nonstoichiometric Titanium Carbide”,

Physica B, 311(3-4), 318-325 (2002) (Crys. Structure, Experimental, *, 18)

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Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 at 25°C [Mas2]

( Ti)

882

2350 at ~20 at.% N

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

a = 295.0 to 297.3

c = 468.2 to 478.8

at 25°C [Mas2]

Ti1-xNx, x = 0.0005-0.22, [V-C2]

(C)

3827(S.P.)

hP4

P63/mmc

C (graphite)

a = 246.12

c = 670.9

at 25°C [Mas2]

’,Ti2C

790 - 770

cF48

Fd3m

Ca33Ge

a = 860 [V-C2]

TiC0.59-TiC0.62 [2002Tas]

ordered phase

’’, Ti2C1+x, (Ti8C5)

770

hR13

R3m or P3121

Ti8C5

a = 611.4 or 305.7

c = 1489.5

[V-C2], [1989Kha], [1991Tas]

ordered phase

TiC

23 (300 K)

hR* a = 294.42 0.03

c = 733.53 0.09

p >18 GPa [1999Dub]

, Ti2N

1080

tP6

P42/mnm

anti-TiO2

a = 494.52

c = 303.42

[V-C2], [1987Len],

[1992Len]

, Ti3N2-x

1103 - 1066

hR6

R3m

AgCrSe2

a = 298.09 0.04

c = 2166.42 0.85

About 30 at.% N

[V-C2], [1986Len2]

[1992Len]

, Ti4N3-x

1291 - 1080

hR8

R3m

Ti7S12

a = 297.95

c = 2896.49

[V-C2], [1986Len1],

[1992Len]

, Ti2N tI16

P41/amd

Ti2N

a = 414.932

c = 878.585

[V-C2]

metastable, [1987Len]

, Ti(CxN1-x)

3290

TiN1-x

TiCx

cF8

Fm3m

NaCl

a (x, T) = 423.13+8.8 x+

(2.338-0.12) x T 10-3 +

(1.0717-0.2258) x T2 10-6

a = 415.9+0.00164x

a = 441.5-0.00348x

a = 430.6 to 432.7

25 - 1200°C, [1994Aig]

(x in Ti(CxN1-x))

x < 0.5 [1987Wri]

x > 0.5

0.51< x < 0.96 [V-C2]

128


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C–N–Ti


Table 3: Coefficients of the Polynomial for the Temperature Diffusivity Data Fit [1995Len]

Table 4: Coefficients of the Polynomial for the Heat Conductivity Data for Stoichiometric Samples


Ti C N

L + + 1820 U L ~96.2

~84

~93.7

~71.5

~2.4

~2.8

~0.7

~12.5

~1.4

~13.2

~5.6

~16

x in Ti(CxN1-x) A B 103 C 107 D 10-4

0.01 8.49 1.84 -7.177 -4.276

0.21 5.61 3.44 -10.94 -0.489

0.40 5.96 1.84 -4.938 -6.508

060 5.52 3.24 -11.53 -9.016

0.79 5.14 3.02 -9.494 1.145

0.99 4.97 3.03 -6.603 -4.805

x in Ti(CxN1-x)0.82

0.00 1.83 3.85 -8.045 0.0602

0.20 1.41 3.94 -8.973 1.809

0.40 1.97 2.97 -5.051 1.922

0.60 2.32 2.20 -3.422 3.210

0.80 2.36 1.79 -2.219 2.698

1.00 2.71 1.33 -1.720 2.694

x in Ti(CxN1-x) A B 103 C 106 D 10-5

0.01 33.453 9.679 -0.147 -8.1666

0.21 19.842 19.090 -3.708 -4.7049

0.40 22.105 11.328 -1.718 -7.6324

060 13.496 29.968 -10.754 -5.4960

0.79 20.098 11.630 -1.143 -5.8377

0.99 15.887 16.740 -2.281 -5.7473

129


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C–N–Ti

20 40 60 80

500

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

Ti C

C, at.%

Te

mp

era

ture

, °C

2276°C

L+δ+(C)

δ+(C)

(C)

δ

L+δ

(βTi)+δ

(αTi)+δ

L

1646.5°C

920.0°C

(βTi)

(αTi)

(αTi)+(βTi)

3066°C

44.53

1670°C

882°C

L+(βTi)

δ´

δ´´

Fig. 1: C-N-Ti.

C-Ti binary phase

diagram

10 20 30 40

250

500

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

3750

4000

Ti 50.00

C 0.00

N 50.00

Ti 50.00

C 50.00

N 0.00C, at.%

Te

mp

era

ture

, °C

L+G

δ+G

L+δ

L

δ+(C)

δ+G+(C)

δ+L+G

3025°C

2776°C

Fig. 2: C-N-Ti.

The TiN-TiC

quasibinary section at

pN2=1 bar

130


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C–N–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti C

N Data / Grid: at.%

Axes: at.%

(αTi)

(βTi)

p1

p2

2900°C

3100

U1 e

1 e1

δ

3300

3400

2900

3100

60

70

80

90

10 20 30 40

10

20

30

40

Ti Ti 50.00

C 50.00

N 0.00

Ti 50.00

C 0.00

N 50.00Data / Grid: at.%

Axes: at.%

δ

ε+δ

(αTi)

(αTi)+δ''

δ''+(αTi)+ε

(αTi)+ε

ε

δ''

δ''+ε

Fig. 3: C-N-Ti.

Liquidus surface

[1996Jon]

Fig. 4: C-N-Ti.


500°C according to

[1978Arb1]

131


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C–N–Ti

60

70

80

90

10 20 30 40

10

20

30

40

Ti Ti 50.00

C 50.00

N 0.00

Ti 50.00

C 0.00


Axes: at.%

L

(βTi)

(αTi)

δ

60

70

80

90

10 20 30 40

10

20

30

40

Ti Ti 50.00

C 50.00

N 0.00

Ti 50.00

C 0.00


Axes: at.%

(βTi)

δ

(αTi)

(βTi)+δ

(αTi)+δ

(αTi)+(βTi)+δ

δ+ζ

(αTi)+ζ(αTi)+δ+ζ

δ+(C)

ζ

Fig. 6: C-N-Ti.


the invariant

temperature 1820°C,

calculated by

[1996Jon]

Fig. 5: C-N-Ti.

Tentative isothermal

section at 1150°C

132


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C–Si–Ti

Carbon – Silicon – Titanium

Yong Du, Honghui Xu, Zhu Pan, Baiyun Huang, Yong Liu, and Huashan Liu

Literature Data

The ternary C-Si-Ti system was investigated by [1965Bru, 1991Wak, 1992Sam, 1995Goe, 1997Nak,

1998Du1, 2000Du]. The literature data up to 1992 were reviewed by [1994McH]. A complete isothermal

section at 1200°C was established by [1965Bru] employing X-ray diffraction (XRD). As many as 70 alloys

were prepared from graphite (1054 ppm total impurities), Si powder (1000 ppm Al, 800 ppm Cu), Ti powder

(1300 ppm C, 1500 ppm H, 1200 ppm Cl), and TiH2 powder (1000-3000 ppm Al, 1000 ppm C, 1500 ppm

H, 1200 ppm Cl) [1965Bru]. A ternary phase with a composition close to Ti-21Si-31C (at.%) was reported

by [1965Bru]. The stoichiometric composition for this ternary phase was shown to be Ti3SiC2 by

subsequent investigations [1967Jei, 1991Wak, 1998Du1, 2000Du]. By means of XRD, polarized light

microscopy, and electron probe microanalysis (EPMA), [1991Wak] determined the isothermal sections at

1250 and 1100°C. In the work of [1991Wak], 19 alloys and 4 diffusion couples were prepared from 99.5

mass% Ti powder, 99.7 mass% Ti rod, technical purity Si powder, single-crystal Si rod, C powder, and

hot-pressed SiC without sintering additives. The results from diffusion couple experiments confirmed the

phase equilibria exhibited by the alloys. [1991Wak] observed a homogeneity range (about 2.5 at.%) for the

ternary phase at 1100°C, but negligible homogeneity range at 1250°C. The results of [1991Wak] agree with

the phase assemblage data from [1965Bru] within estimated experimental errors. The phase relationships at

high pressures from 10 to 20 kbar were investigated by [1992Sam] using a diffusion couple technique. The

diffusion couples consisted of solid cylindrical rods of SiC encapsulated inside Ti cans. They were annealed

at high temperatures and pressures in a piston and cylinder device. The quenched interface was

characterized by backscattered electron imaging and EPMA. The work of [1992Sam] indicates that Ti3Si is

stabilized at pressures greater than 2.4 kbar exhibiting an appreciable solubility for C (up to 9 at.%) at

1200°C. At higher temperatures and lower pressures, it decomposes to Ti5Si3, TiC, and TiSi. A schematic

isothermal section at 1200°C for pressures greater than approximately 2.4 kbar was presented by

[1992Sam]. This schematic isothermal section is nearly the same as that at 1 bar published by [1965Bru]

except for the introduction of phase equilibria involving Ti3Si. More recently, [1998Du1, 2000Du]

determined the isothermal section at 1200°C and the isopleths at 5, 10, and 15 at.% C. The authors of

[1998Du1, 2000Du] prepared 21 alloys starting from powdered C (spectroscopically pure), Si

(99.9 mass%), and Ti (99 mass%). DTA using a heating and cooling rate of 5°C/min was the primary

experimental method, supplemented by XRD examinations of the alloys. In addition, the transition

temperature for L+TiC Ti3SiC2+SiC at 2200 20°C was measured by [2000Du] using the Pirani

technique. By considering reliable literature data [1965Bru, 1991Wak] and their own experimental data,

[1998Du1, 2000Du] obtained an optimized set of thermodynamic parameters for the whole ternary C-Si-Ti

system. The calculations of [1998Du1, 2000Du] satisfactorily account for most of the experimental data.

From the observation of diffusion paths between SiC and Ti, [1995Goe] and [1997Nak] constructed

isothermal sections from 700 to 1200°C and at 1400°C, respectively. These proposed sections are

essentially based on the measurements of [1965Bru].

By XRD examination of deposits obtained by gas-phase crystallization, [1975Aiv] found that the solubility

of SiC in TiC reaches 11.1 mass% in the temperature range of 1100 to 1700°C. By comparing the Gibbs

energy changes of all possible reactions during the crystallization, however, [1975Aiv] suggested that some

other reactions could occur as well as the dominant reaction leading to the formation of TiC. Consequently,

the equilibrium solubility of SiC in TiC is less than 11.1 mass%. Such a high solubility was confirmed by

[1998Wit]. Three groups of authors demonstrated via extensive measurements [1965Bru, 1991Wak] and

thermodynamic modeling [2000Du] that the solubility of TiC in SiC is negligible.

The enthalpy of formation of Ti3SiC2 was estimated by [2002Kis] using a version of DTA that used the

lattice spacing of TiCx as an indication of temperature.

133


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C–Si–Ti

Several preliminary thermodynamic calculations for the ternary C-Si-Ti system have been performed

[1956Bre, 1994Rac2, 1994Sei, 1998Zim]. Based on estimated thermodynamic properties for binary phases

only, [1956Bre] constructed the isothermal section at 1727°C. In view of the simplifying and often

unrealistic assumptions in the calculations [1994Rac2, 1994Sei, 1998Zim], these results are not reported in

this evaluation.

Data relating to the crystal structure of the Ti3SiC2 ternary phase have been reported by [1967Jei, 1972Nic,

1987Got, 1995Aru, 1998Kis, 1999Gam, 1999Rad, 2002Tan2, 2002Yu]. This ternary phase was prepared

via solid-state reaction [1967Jei], chemical vapor deposition (CVD) [1972Nic, 1987Got], and an arc

melting and annealing route [1995Aru, 2000Du]. Other techniques used to synthesize Ti3SiC2 include hot

isostatic pressing (HIP), HP (hot pressing), combustion synthesis, electron beam ignited solid state reaction,

and self propagating high temperature synthesis (SHS).

[2000Ono] studied the pressure-volume dependence of Ti3SiC2 up to 60 GPa (600 kbar) under static

conditions. [2003Jor] extended these data for ultra high pressures of up to 120 GPa (1200 kbar) under

dynamic (shock loading) conditions. An indication of a possible unspecified phase transition was obtained

at 90 to 120 GPa.

Binary Systems

The C-Si phase diagram is accepted from [1996Gro], and is presented in Fig. 1. The C-Ti system has been

assessed by [1996Jon, 1996Sei1, 1998Sun]. The evaluation by [1998Sun] is adopted since it gives a better

overall agreement with the experimental data in some ternary systems, as shown by [1998Du2, 2000Du].

Figure 2 shows the C-Ti phase diagram. The Si-Ti system was reevaluated by [1998Sei], who made a

noticeable improvement of the previous modeling [1996Sei2] by considering the newly measured phase

diagram data [1998Du3]. The Si-Ti phase diagram is reproduced in Fig. 3.

Solid Phases

Table 1 summarizes the crystal structure data for all phases in the ternary C-Si-Ti system. Based on EPMA

measurements on alloys annealed for 100 h at 1200°C, [1995Aru] proposed that Ti3SiC2 exists over a range

of compositions at this temperature. The finding of [1995Aru] appears to be in some disagreement with the

results of [1991Wak, 1992Sam, 2000Du]. Employing EPMA, both groups of authors [1991Wak, 1992Sam]

reported small homogeneity ranges for Ti3SiC2 in the temperature range 1000 to 1250°C. The results of

[1991Wak] are preferable because of the extended period of annealing (316 h at 1000 and 1250°C).

Using high-resolution electron microscopy (HREM), [2002Yu] discovered a modification of Ti3SiC2,

differing from by its stacking sequence. It was obtained together with ordinary Ti3SiC2 following hot

isostatic pressing. Owing to the lack of detailed structure data, it could not be included in Table 1.


[1998Du1, 2000Du] measured six invariant ternary reaction temperatures by means of DTA and Pirani

techniques. They were used in their thermodynamic calculation of the system. The computed reaction

temperatures and the compositions of the respective phases are given in Table 2. The calculated reaction

temperatures are in good agreement with the experimental data [1998Du1, 2000Du]. The complete reaction

scheme [2000Du] is presented in Fig. 4. It follows the thermodynamic modeling by [2000Du], which is in

agreement with the selected experimental data [1965Bru, 1991Wak, 2000Du]. The calculated invariant

reactions in the solid state are predicted from the modeling.

Liquidus Surface

The computed liquidus surface in Fig. 5a is taken from the thermodynamic modeling by [2000Du]. Three

degenerate reactions (D1, D2, and D3) are shown in Fig. 5b, where the liquidus surface close to the Si-Ti

binary side is presented.

134


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C–Si–Ti

Isothermal Sections

The isothermal sections in Figs. 6 to 8 are taken from the thermodynamic calculation of [2000Du]. The

shape of the extension of Ti5Si3Cx into the ternary system shown in Figs. 7 and 8 was amended slightly to

bring it into agreement with the experimental data of [1991Wak]. Such a correction, most probably, should

be also carried out in Fig. 6, but there are no experimental data for that temperature, therefore a dashed line

shows it as uncertain.


Figures 9 to 11 present the calculated isopleths at 5, 10, and 15 at.% C, respectively [2000Du]. The

calculated liquidus below 1500°C was confirmed by DTA measurements using selected alloys [2000Du].

Thermodynamics

The thermodynamic model parameters of the phases as obtained by [2000Du] are listed in that work.

The Gibbs energy of formation of Ti3SiC2 ( fG°) relative to the solid forms of the elements has been

estimated by several groups of researchers [1990Sam, 1994Rac2, 1994Sei, 2000Du]. [1990Sam] derived

fG° from (i) measurement of the mole fraction of C in TiCx in equilibrium with SiC and Ti3SiC2 in the

temperature range of 1200 to 1500°C, and (ii) the thermodynamic models of [1984Tey, 1984Urh] used for

TiCx. For both models [1984Tey, 1984Urh], the resulting fG° decreases with temperature. Using the

limited experimental phase relationships [1965Bru, 1991Wak], [1994Sei] found a similar behavior for

fG°. However, [1994Rac2] found that fG° increases with temperature considering the experimental

results of [1990Sam] and the sublattice model proposed for TiCx [1989Vin]. The finding of [1994Rac2] was

confirmed by [2000Du], who optimized fG° using experimental phase diagram data covering wide

temperature and composition ranges. The negative entropy of formation results in an increase in fG° with

temperature, which is consistent with the thermal stability of Ti3SiC2 [1994Rac1]. The negative entropy of

formation is also expected from a solid phase with a strongly negative enthalpy of formation. The fG°

value given by [2000Du] is - 91191+4.14083 T kJ (mol-atoms)-1. The fH value of -76 kJ (mol-atoms)-1

for this phase was estimated experimentally by [2002Kis] using “diffraction DTA” (see above for details).

The present authors considered their result as being too low, based on a comparison with other phases in the

system.

[1992Sam] found that Ti3Si was stabilized with C at high pressures and estimated its Gibbs energy of

formation for a saturated solution with the composition Ti-22.7Si-9C (at.%) at 1300°C and 14 kbar through

an evaluation of the experimental phase diagram and the available thermodynamic data for the

corresponding binary phases.

Using first-principle technique, [2000Wil3] calculated the enthalpies of formation of Ti5Si3Cx with two

compositions Ti-36.4Si-3C and Ti-35.3Si-5.9C (at.%). The enthalpies of formation obtained are -76775

J (mol-atoms)-1 for Ti-36.4Si-3C and -84029 J (mol-atoms)-1 for Ti-35.3Si-5.9C. [2000Wil3] also

computed the enthalpy of formation of the Ti5Si3 phase. The calculated enthalpy is in good agreement with

the assessed value [1996Sei2].

The low temperature heat capacity of Ti3SiC2 was measured between 2 and 10 K by [1999Ho1, 1999Ho2]

using a standard adiabatic calorimeter in a liquid helium cryostat. The high temperature heat capacity of this

ternary compound was measured from 0 to 1023°C by [1999Bar].


Ti3SiC2 is a novel structural/functional material due to its unique combination of materials properties, and

its synthesis and properties have been the subject of much study [1959Now, 1970Nic, 1980Cha, 1981Pai,

1984Mar, 1985Mor, 1988Bor, 1989Bac, 1989Pam, 1991Jer, 1993Got, 1994Aru, 1994Rac1, 1996Wen,

1997Via, 1997Gol, 1997Kel, 1998Goe, 1998Kis, 1999ElR, 2000Bar1, 2000Bar2, 2000Che, 2000Fin,

2000Tho, 2000Wil1, 2000Wil2, 2000Zho1, 2000Zho2, 2001Tan, 2002Gao, 2002Kis, 2002Liu, 2002Mam,

2002Ril, 2002Sun, 2002Tan1, 2002Yok, 2002Yu]. This compound has a low density (4.53 g cm-3), a high

melting temperature (> 2300°C), high modulus, good thermal and electrical conductivity, excellent thermal

135


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C–Si–Ti

shock resistance and high-temperature oxidation resistance because of the protective coating of TiO2 and

SiO2 formed in air. The potential applications of Ti3SiC2 are myriad and far-reaching.

Ti3SiC2 can be widely used for high-temperature structural applications because of its high Young’s and

shear moduli, chemical stability and covalent bonding. The density of Ti3SiC2 is roughly half the density

of current Ni-based superalloys, and it possesses reasonable mechanical properties at temperatures that

would render the best superalloys on the market today unusable. Ti3SiC2 behaves like a “soft” ceramic

material making it easily machinable. Low density Ti based materials reinforced with SiC fibers are used

in microelectronic devices, metal-ceramic joints and metal matrix composites.

The main advantage of using Ti3SiC2 is that it is easily machinable in its final fired state. Conventional

ceramics require a sintering step after machining which results in more than 2 % shrinkage.

Ti3SiC2 can be used as kiln furniture mainly because of its excellent oxidation resistance, ease of

machinability, relatively low cost of raw materials, and excellent thermal shock and chemical resistance.

Ti doped with SiC via ion implantation or excimer laser surface processing can be used for tribological

applications because these processes modify the microstructure and chemistry of the surface and thus

reduce friction and wear of the materials. Carburizing Ti3SiC2 can increase the surface hardness. The treated

surfaces are wear and corrosion resistant.

Ti3SiC2 can be used in heat exchangers. Ti3SiC2 is an excellent thermal conductor with a conductivity that

does not decrease significantly with increasing temperature.

Miscellaneous

[1992Sam] studied the temperature vs. pressure region of stability of Ti3SiC0.4 (solid solution of C in Ti3Si

compound) with respect to the reaction Ti3Si(C) TiCx+Ti5Si3+Ti(Si,C). These data are presented in

Fig. 12.

Two groups of authors [2000Tho, 2000Wil1] have investigated the effect of C additions on the structure of

Ti5Si3 by means of XRD and neutron diffraction. Figure 13 shows the variation in the lattice parameter of

Ti5Si3Cx with respect to C content.

[1987Got] prepared Ti3SiC2 polycrystalline plates by means of chemical vapor deposition (CVD) using

SiCl4, TiCl4, CCl4, and H2 as source gases. The CVD phase diagram of [1987Got] is shown in Fig. 14.

Figure 15 presents the pressure dependence of the reduced volume (V/V0) as obtained by [2000Ono] and

[2003Jor] for ultra-high pressures (up to 120 GPa). The small deviation of the experimental points of

[2003Jor] from the extrapolated results of [2000Ono] was treated as indication of a phase transition, though

this evidence seems to be too weak to be accepted here.

Further experimental work is necessary to confirm the invariant reaction temperatures in the solid state. It

would also be of interest to investigate whether (i) Ti3SiC2 exhibits a homogeneity range and (ii) whether

TiC dissolves noticeable amounts of SiC.

References

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J. Electrochem. Soc., 103(1), 38-51 (1956) (Calculation, Equi. Diagram, 52)

[1959Now] Nowotny, H., Kieffer, R., Benesovsky, F., Laube, F., “Carbides, Silicides, and Borides of

High Melting Point” (in German), Acta Chimi. Acad. Sci. Hung., 18, 35-44 (1959)


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136


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C–Si–Ti

[1972Nic] Nickl, J.J., Schweitzer, K.K., Luxenberg, P., “Gas Phase Deposition in the Ti-Si-C System”

(in German), J. Less-Common Met., 26, 335-353 (1972) (Experimental, Crys. Structure,

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72(2), 305-311 (1980) (Experimental, Kinetics, 10)

[1981Pai] Pailler, R., Martineau, P., Lahaye, M., Naslain, R., “Fiber-Matrix Chemical Interactions at

High Temperatures and Composite Strength in B (or SiC) Filament-Ti Base Matrix

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Ti3SiC2”, J. Eur. Ceram. Soc., 5(5), 283-287 (1989) (Experimental, Crys. Structure, 14)

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Stereochemistry of Titanium Carbide Deposited Between a Graphite Substrate and a

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81-96 (1989) (Experimental, Thermodyn., 19)

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[1991Jer] Jervis, T.R., Hirvonen, J.-P., Nastasi, M., “Tribology and Mechanical Properties of Excimer

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Pressures”, J. Mater. Res., 7(6), 1473-1479 (1992) (Experimental, Equi. Diagram,

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J. Mater. Res., 8(10), 2725-2733 (1993) (Experimental, Kinetics, 16)

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[1994McH] McHale, A.E., “V. Silicon Plus Carbon Plus Metal(s) (a) Si Plus C- Plus Metal”, Phase

Equilibria Diagrams, Phase Diagrams for Ceramists, 10, 16-19 (1994) (Equi. Diagram,

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Ternary Phase Ti3SiC2”, J. Mater. Sci., 29, 3384-3392 (1994) (Experimental, Kinetics, 18)

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from TiCl4-SiCl4-CH4-H2 Gas Mixtures: Part I A Thermodynamic Approach”, J. Mater.

Sci., 29, 5023-5040 (1994) (Calculation, Thermodyn., 18)

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Minerals Ressources (RAMM’94), 1994, Penang, Malaysia, Perak Foundation, Ipoh, ISBN

938-9700-24-3, 11-20 (1994) (Calculation, Equi. Diagram, 38)

[1995Aru] Arunajatesan, S., Carim, A.H., “Synthesis of Titanium Silicon Carbide”, J. Am. Ceram.

Soc., 78(3), 667-672 (1995) (Experimental, Equi. Diagram, 16)

[1995Goe] Goesmann, F., Schmid-Fetzer, R., “Temperature-Dependence, Interface Reactions and

Electrical Contact Properties of Titanium on 6H-SiC”, Semicond. Sci. Technol., 10,


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System Al-Si-C”, Calphad, 20, 247-254 (1996) (Calculation, Equi. Diagram, #, 37)

[1996Jon] Jonsson, S., “Assessment of the Ti-C System”, Z. Metallkd., 87, 703-712 (1996)

(Calculation, Equi. Diagram, 59)

[1996Sei1] Seifert, H.J., Lukas, H.L., Petzow, G., “Thermodynamic Optimization of the Ti-C System”,

J. Phase Equilib., 17, 24-35 (1996) (Calculation, Equi. Diagram, 58)

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Z. Metallkd., 87(1), 2-13 (1996) (Calculation, Equi. Diagram, 63)

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Synthesized TiC and SiC”, Int. J. Powder Met., 32(3), 259-264 (1996) (Experimental,

Mechan. Prop., 10)

[1997Nak] Naka, M., Feng, J.C., Schuster, J.C., “Phase Reaction and Diffusion Path of the SiC/Ti

System”, Metall. Mater. Trans. A, 28A(6), 1385-1390 (1997) (Experimental, Equi.

Diagram, 11)

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Al-C-Si-Ti Quaternary System: an Experimental Approach”, Mater. Sci. Eng. A, 229,

95-113 (1997) (Equi. Diagram, Review, 35)

[1997Gol] Goldin, B.A., Istomin, P.V., Ryabkov, Yu.I., “Reduction Solid-State Synthesis of Titanium

Silicide Carbide, Ti3SiC2”, Inorg. Mater., 33(6), 577-579 (1997), translated from Neorg.

Mater., 33(6) 691-693 (1997) (Experimental, Crys. Structure, 9)

[1997Kel] Kellerman, D.G., Gorshkov, V.S., Blinovskov, Ya.N., Grigorov, I.G., Perelyaev, V.A.,

Shveikin, G.P., “Synthesis and Properties of the Ternary Phase Ti3SiC2”, Inorg. Mater.

33(3), 271-274 (1997), translated from Neorg. Mater., 33(3), 329-332 (1997)


[1998Du1] Du, Y., Schuster, J.C., “Experimental and Thermodynamic Investigations in the Ti-Si-C

System”, Ber. Bunsen-Ges. Phys. Chem., 102(9), 1185-1188 (1998) (Experimental,

Calculation, Equi. Diagram, Thermodyn., 25)

[1998Du2] Du, Y., Schuster, J.C., “Experimental Investigation and Thermodynamic Modeling of the

Ni-Ti-C System”, Z. Metallkd., 89(6), 399-410 (1998) (Calculation, Experimental, Equi.

Diagram, Thermodyn., 39)

[1998Du3] Du, Y., Schuster, J.C., “A Reinvestigation of the Constituent of the Partial System TiSi-Si”,

J. Mater. Sci. Lett., 17, 1407-1408 (1998) (Experimental, Equi. Diagram, 7)

[1998Goe] Goesmann, F., Wenzel, R., Schmid-Fetzer, R., “Preparation of Ti3SiC2 by

Electron-Beam-Ignited Solid-State-Reaction”, J. Am. Ceram. Soc., 81(11), 3025-3028

(1998) (Experimental, Electr. Prop., Mechan. Prop., 10)

138


MSIT®

C–Si–Ti

[1998Kis] Kisi, E.H., Crossley, J.A.A., Myhra, S., Barsoum, M.W., “Structure and Crystal Chemistry

of Ti3SiC2”, J. Phys. Chem. Solids, 59(9), 1437-1443 (1998) (Experimental, Crys.

Structure, 19)

[1998Sei] Seifert, H.J., Unpublished results, Max-Planck-Institut fuer Metallforschung, Stuttgart,

Germany (1998) (Calculation, Equi. Diagram, #) as quoted by [2000Du]

[1998Sun] Sundman, B., Dumitrescu, L.F.S., “Extrapolations Based on Ti-C-N”, in “COST

507-Definition of Thermochemical and Thermophysical Properties to Provide a Database

for the Development of New Light Alloys”, Proceedings of the Final Workshop,Vol. 1,

COST Secretariat, Brussels, Belgium, 173-206 (1998) (Calculation, Equi. Diagram, #, 60)

[1998Wit] Witthaut, M., Weiss, R., Zimmermann, E., von Richthofen, A., Neuschutz, D., “The

Formation of Interface Phases in the Diffusion Couple Ti-SiC”, Z. Metallkd., 89(9), 623-628


[1998Zim] Zimmermann, E., Weiss, R., Witthaut, M., Neuschuetz, D., “Interdiffusion Paths in the

Diffusion Couple Ti-SiC”, Z. Metallkd., 89(10), 714-718 (1998) (Calculation, Equi.

Diagram, Thermodyn., 18)

[1999Bar] Barsoum, M.W., El-Raghy, T., Rawn, C.J., Porter, W.D., Wang, H., Payzant, A., Hubbard,

C., “Thermal Properties of Ti3SiC2”, J. Phys. Chem. Solids, 60, 429-439 (1999)


[1999ElR] El-Raghy, T., Barsoum, M.W., “Processing and Mechanical Properties of Ti3SiC2: I,

Reaction Path and Microstructure Evolution”, J. Am. Ceram. Soc., 82(10), 2849-2854


[1999Gam] Gamarnik, M.Y., Barsoum, M.W., “Bond Lengths in the Ternary Compounds Ti3SiC2,

Ti3GeC2 and Ti2GeC”, J. Mater. Sci., 34, 169-174 (1999) (Experimental, Crys.

Structure, 15)

[1999Ho1] Ho, J.C., Hamdeh, H.H., Barsoum, M.W., El-Raghy, T., “Low Temperature Heat Capacity

of Ti3SiC2”, J. Appl. Phys., 85(11), 7970-7971 (1999) (Experimental, Thermodyn., 14)

[1999Ho2] Ho, J.C., Hamdeh, H.H., Barsoum, M.W., El-Raghy, T., “Low Temperature Heat Capacities

of Ti3Al1.1C1.8, Ti4AlN3, and Ti3SiC2”, J. Appl. Phys., 86(7), 3609-3611 (1999)


[1999Rad] Radhakrishnan, R., Williams, J.J., Akinc, M., “Synthesis and High-Temperature Stability of

Ti3SiC2”, J. Alloys Compd., 285, 85-88 (1999) (Experimental, Crys. Structure, 11)

[2000Bar1] Barsoum, M.W., El-Raghy, T., Radovic, M., “Ti3SiC2: A Layered Machinable Ductile

Carbide”, Interceram, 49(4), 226-233 (2000) (Crys. Structure, Thermodyn., Mechan. Prop.,

Electr. Prop., Thermal Conduct., Corrosion, Review, 56)

[2000Bar2] Barsoum, M.W., “The Mn+1AXn Phases: A New Class of Solids; Thermodynamically

Stable Nanolaminates”, Prog. Solid St. Chem., 28, 201-281 (2000) (Crys. Structure,

Mechan. Prop., Electr. Prop., Thermal Conduct., Corrosion, Review, 209)

[2000Che] Chen, H., Yin, Y.Z., He, Y.J., “Thermoelectric Properties of Ti-Doped SiC/TiCx Composite

Synthesized by TPPP Method”, Modern Phys. Lett. B, 14(4), 131-138 (2000)

(Experimental, Phys. Prop., Thermal Conduct., 9)

[2000Du] Du, Y., Schuster, J.C., Seifert, H.J., Aldinger, F., “Experimental Investigation and

Thermodynamic Calculation of the Titanium-Silicon-Carbon System”, J. Am. Ceram. Soc.,

83(1), 197-203 (2000) (Experimental, Calculation, Crys. Structure, Equi. Diagram,

Thermodyn., #, *, 41)

[2000Fin] Finkel, P., Barsoum, M.W., El-Raghy, T., “Low Temperature Dependencies of the Elastic

Properties of Ti4AlN3, Ti3Al1.1C1.8, and Ti3SiC2”, J. Appl. Phys., 87(4), 1701-1703 (2000)


[2000Ono] Onodera, A., Hirano, H., Yuasa, T., Guo, N.F., Miyaamoto, G., in “Science and Technology

of High Pressure”, Mangnani M.H., Nellis W.J., Nicol M.F. (Eds.), Univ. Press, India, 2000

(Experimental, Equi. Diagram, #)

[2000Tho] Thom, A.J., Young, V.G., Akinc, M., “Lattice Trends in Ti5Si3Zx (Z = B, C, N, O and

0 < x < 1)”, J. Alloys Compd., 296, 59-66 (2000) (Experimental, Crys. Structure, 32)

139


MSIT®

C–Si–Ti

[2000Wil1] Williams, J.J., Kramer, M.J., Akinc, M., Malik, S.K., “Effects of Interstitial Additions on

the Structure of Ti5Si3”, J. Mater. Res., 15(8), 1773-1779 (2000) (Experimental, Crys.

Structure, 16)

[2000Wil2] Williams, J.J., Kramer, M.J., Akinc, S.K., “Thermal Expansion of Ti5Si3 with Ge, B, C, N,

or O Additions”, J. Mater. Res., 15(8), 1780-1785 (2000) (Experimental, Phys. Prop., 15)

[2000Wil3] Williams, J.J., Ye, Y.Y., Kramer, M.J., Ho, K.M., Hong, L., Fu, C.L., Malik, S.K.,

“Theoretical Calculations and Experimental Measurements of the Structure of Ti5Si3 with

Interstitial Additions”, Intermetallics, 8, 937-943 (2000) (Theory, Experimental, Crys.

Structure, Thermodyn., 22)

[2000Zho1] Zhou, Y.C., Sun, Z.M., Sun, J.H., Zhang, Y., Zhou, J., “Titanium Silicon Carbide: a

Ceramic or a Metal?”, Z. Metallkd., 91(4), 329-334 (2000) (Crys. Structure, Mechan. Prop.,

Phys. Prop., Review, 19)

[2000Zho2] Zhou, Y.C., Sun, Z.M., Yu, B.H., “Microstructure of Ti3SiC2 Prepared by the In-Situ Hot

Pressing/Solid-Liquid Reactrion Process”, Z. Metallkd., 91(11), 937-941 (2000)


[2001Tan] Tang. K., Wang, C.A., Huang, Y., Zan, Q.F., “Growth Model and Morphology of Ti3SiC2

Grains”, J. Cryst. Growth, 222, 130-134 (2001) (Experimental, Crys. Structure, 11)

[2002Gao] Gao, N.F., Miyamoto, Y., Zhang, D., “On Physical and Thermochemical Properties of

High-Purity Ti3SiC2”, Mater. Lett., 55, 61-66 (2002) (Experimental, Phys. Prop.,

Thermochem. Prop., 21)

[2002Kis] Kisi, E.H., Riley, D.P., “Diffraction Thermometry and Differential Thermal Analysis”, J.

Appl. Crystallogr., 36(6), 664-668 (2002) (Crys. Structure, Experimental, Thermodyn., 14)

[2002Liu] Liu, G.M., Li, M.S., Zhou, Y.C., “Corrosion Behavior of Ti3SiC2 and Siliconized Ti3SiC2

in the Mixture of K2CO3 and Li2CO3 Melts at 750°C”, J. Mater. Sci. Lett., 21(22),

1755-1757 (2002) (Experimental, Corrosion, 9)

[2002Mam] Mamyan, S.S., “Thermodynamic Analysis of the SHS Processes”, Key Eng. Mater., 217,

1-8 (2002) (Calculation, Thermodyn., 16)

[2002Ril] Riley, D.P., Kisi, E.H., Hansen, T.C., Hewat, A.W., “Self-Propagating High-Temperature

Synthesis of Ti3SiC2: I, Ultra-High-Speed Neutron Diffraction Study of the Reaction

Mechanism”, J. Am. Ceram. Soc., 85(10), 2417-2424 (2002) (Experimental, Crys.

Structure, 48)

[2002Sun] Sun, Z.M, Zhou, Y.C., “High Temperature Oxidation and Hot Corrosion Bahavior of

Ti3SiC2-Based Materials”, Key Eng. Mater., 224-226, 791-796 (2002) (Experimental,

Corrosion, 14)

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Mechanism and Growth of Ti3SiC2 Synthesized by Hot-Pressing”, Mater. Sci. Eng. A, 328,

206-212 (2002) (Theory, Crys. Structure, 22)

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Ti3SiC2”, Mater. Lett., 55, 50-55 (2002) (Experimental, Crys. Structure, 12)

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Windows with Gem and CHD”, Calphad, 26(2), 155-166 (2002) (Calculation, Equi.

Diagram, 9)

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Res, 17(5), 948-950 (2002) (Experimental, Crys. Structure, 18)

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“High Pressure Behavior of Titanium-Silicon Carbide (Ti3SiC2)”, J. Appl. Phys., 93,

9639-9643 (2003) (Experimental, Equi. Diagram, #, 8)

140


MSIT®

C–Si–Ti


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(C)

< 3827 (S.P.)

hP4

P63/mmc

C (graphite)

a = 246.12

c = 670.9

25°C [Mas2]

(Si)

< 1414

cF8

Fd3m

C (diamond)

a = 543.06 25°C [Mas2]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

25°C [Mas2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 25°C [Mas2]

Ti3Si

< 1170

tP32

P42/n

Ti3P

a = 1019.6

c = 509.7

[V-C]

Ti5Si3Cx (T2)

< 2130

hP16

P63/mcm

Mn5Si3

a = 742.9

c = 513.92

a = 745.2

c = 514.8

a = 745.43 0.02

c = 514.74 0.06

a = 746 2

c = 515.2 0.2

at x = 0 [V-C]

at x = 0 [1970Nic]

at x = 0 [2000Tho]

at x = 0 [2000Wil1]

Ti5Si4< 1920

tP36

P41212

Zr5Si4

a = 670.2

c = 1217.4

[V-C]

TiSi

< 1570

oP8

Pmm2

TiSi

a = 361.8

b = 649.2

c = 497.3

[V-C]

TiSi2< 1489

oF24

Fddd

TiSi2

a = 826.7

b = 480.0

c = 855.0

[V-C]

TiCx

< 3067

cF8

Fm3m

NaCl

a = 430.6 to 432.7 0.51 < x < 0.96 [V-C2]

141


MSIT®

C–Si–Ti


SiC

< 2823

cF8

F43m

ZnS

a = 435.8 [1996Gro]

* 1, Ti3SiC2

< 2374

hP12

P63/mmc

MoC

a = 306.8 0.2

c = 1766.9 0.6

a = 306.6

c = 1764.6

a = 306.4

c = 1765.0

a = 307

c = 1769

a = 305.75 0.01

c = 1762.35 0.03

a = 306.945 0.003

c = 1767.749 0.01

a = 306.65 0.05

c = 1766.9 0.3

a = 306.557 0.006

c = 1763,00 0.05

a = 307.378 0.008

c = 1768,03 0.07

a = 307.827 0.007

c = 1771,19 0.06

a = 308.314 0.008

c = 1774.23 0.07

a = 308.896 0.007

c = 1777.84 0.05

[1967Jei]

[1972Nic]

[1987Got]

[1994Aru]

[1998Kis]

[1999Rad]

[1999Gam]

at 25°C [1999Bar]

at 355°C [1999Bar]

at 531°C [1999Bar]

at 714°C [1999Bar]

at 906°C [1999Bar]


C Si Ti

L + TiCx SiC + 1 2203 U3 L

TiCx

SiC

1

15.66

48.37

50

33.33

54.26

0.185

50

16.67

30.08

51.45

0

50

L TiSi + TiSi2, Ti5Si3Cx 1485.3 D2 L

TiSi

TiSi2Ti5Si3Cx

6.92 10-3

0

0

10.97

64.58

50

66.67

33.39

35.41

50

33.33

55.64

L + Ti5Si3Cx 1 + TiSi2 1482.6 U4 L

Ti5Si3Cx

1

TiSi2

8.82 10-2

11.10

33.33

0

69.48

33.34

16.67

66.67

30.43

55.56

50

33.33

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

142


MSIT®

C–Si–Ti

L + 1 TiSi2 + SiC 1472.6 U5 L

1

TiSi2SiC

0.109

33.33

0

50

71.47

16.67

66.67

50

28.42

50

33.33

0

L + TiCx ( Ti) + Ti5Si3Cx 1353.5 U6 L

TiCx

( Ti)

Ti5Si3Cx

0.02

34.17

0.798

2.49

14.73

6.15 10-9

3.60

32.51

85.25

65.83

95.59

64.99

L (Si) + TiSi2, SiC 1330.1 D3 L

(Si)

TiSi2SiC

4.79 10-3

5.95 10-4

0

50

85.82

99.99

66.67

50

14.17

8.06 10-9

33.33

0


C Si Ti

20 40 60 80

0

250

500

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

3750

4000

4250

4500

4750

5000

Si C

C, at.%

Te

mp

era

ture

, °C

2823

1413.84

L+(C)

SiC+(C)

L

(Si)+SiC

L+SiC

SiC

Fig. 1: C-Si-Ti.

The C-Si binary

system [1996Gro]

143


MSIT®

C–Si–Ti

20 40 60 80

0

250

500

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

3750

4000

4250

4500

4750

5000

Ti C

C, at.%

Te

mp

era

ture

, °C

2772

1653

919

L

TiCx

(αTi)+TiCx

(βTi)+TiCx

L+TiCx

L+(C)

SiC+(C)

Fig. 2: C-Si-Ti.

The C-Ti binary

system [1998Sun]

20 40 60 80

0

250

500

750

1000

1250

1500

1750

2000

2250

2500

Ti Si

Si, at.%

Te

mp

era

ture

, °C

L

862

Ti3Si

1170

1337

(αTi)

(βTi)

2130

1920

1570

TiSi2

TiSi2+(Si)

1485

TiSi

Ti5Si4

TiSi+TiSi2

1300

1489

Ti5Si3

(αTi)+Ti3Si

(βTi)+Ti3Si

Fig. 3: C-Si-Ti.

The Si-Ti binary

system [1998Sei]

144


MSIT®

C–Si–Ti

Fig

. 4:

C

-Si-

Ti.

Rea

ctio

n s

chem

e [2

000D

u]

C-T

iS

i-T

iC

-Si

C-S

i-T

i

l +

Τ2

Ti 5

Si 4

19

20

p4

l T

iCx +

(C

)

27

72

e 1

L +

(C

) T

iCx +

SiC

23

58

U1

l +

(C

) S

iC

28

23

p1

L +

TiCx

Τ 2

24

06

p2m

ax

l T

iCx +

(βT

i)

16

53

e 2

(βT

i) +

TiCx

(αT

i)

91

9p

7

l +

Ti 5

Si 4

TiS

i

15

70

p5

l T

iSi

+ T

iSi 2

14

85

e 4

l (

βTi)

+ T

2

13

37

e 6

l T

iSi 2

+ (

Si)

13

30

e 7

(βT

i) +

T2

Ti 3

Si

11

70

p6

(βT

i)(α

Ti)

+ T

i 3S

i

86

2e 8

l (

Si)

+ S

iC

14

14

e 5

L +

TiCx

τ 1

23

74

p3m

ax

L +

TiCx

T2 +

τ1

23

34

U2

L +

TiCx

SiC

+ τ

12

203

U3

L +

Ti 5

Si 4

TiS

i, T

21

570

D1

L T

iSi

+ T

iSi 2

, T

21

485

D2

LΤ 2

+ T

iSi 2

14

89

e 3m

ax

L +

T2

TiS

i 2 +

τ1

14

83

U4

L +

τ1

TiS

i 2 +

SiC

14

73

U5

L +

ΤiCx

(βT

i) +

T2

13

54

U6

L T

iSi 2

+ (

Si)

, S

iC1

330

D3

(βT

i) +

T2

TiCx+

Ti 3

Si

10

95

U7

SiC

+ T

iCx

(C

) +

τ1

90

6U

8

(βT

i) +

TiCx

(αT

i) +

Ti 3

Si

88

4U

9

L+

TiCx+

SiC

(C)+

TiCx+

SiC

TiCx+

T2+

τ 1L

+T

2+

τ 1

TiCx+

SiC

+τ 1

L+

SiC

+τ 1

L+

TiS

i+T

2T

i 5S

i 4+

TiS

i+T

2

TiS

i+T

iSi 2

+T

2

T2+

TiS

i+τ 1

L+

TiS

i 2+

SiC

TiS

i 2+

SiC

+τ 1

TiCx+

(βT

i)+

T2

TiS

i 2+

(Si)

+S

iC

(βT

i)+

TiCx+

Ti 3

Si

T2+

TiCx+

Ti 3

Si

SiC

+(C

)+τ 1

TiCx+

(C)+

τ 1

(βT

i)+

(αT

i)+

Ti 3

Si

TiCx+

(αT

i)+

Ti 3

Si

145


MSIT®

C–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti C


Axes: at.%

SiC

τ1

U3

U1

(C)

TiCx

T2

U2

3600

3000

25502350

e1

3000

2800

2550

23501900e2

1900

p3max

p2max

p1

2800

(Ti)

D3 SiC

T2

TiSi2

D2TiSi

Ti Si5 4

D1

100

e7

80

e4

60

p5

p4

40

e ,max3U4

τ1

20

0.00.0 0.02 0.04 0.06 0.08 0.1

TiC

U6 ( Ti)β

e6

Si,

at.

%

C, at.%

(Si)

Fig. 5a: C-Si-Ti.

The liquidus surface

[2000Du]

Fig. 5b: C-Si-Ti.

Enlarged part of

Fig. 5(a) in the

composition range 0

to 0.1 at.% C

146


MSIT®

C–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti C


Axes: at.%

(βTi)

τ1

SiC

TiSi2

TiSi

Ti5Si

4

T2

TiCx

20

40

60

80

20 40 60 80

20

40

60

80

Ti C


Axes: at.%

(βTi)

τ1

SiC

TiSi2

TiSi

Ti5Si

4

T2

TiCx

(Si)

L

L

Fig. 7: C-Si-Ti.


1250°C [2000Du]

Fig. 6: C-Si-Ti.


1400°C [2000Du]

147


MSIT®

C–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti C


Axes: at.%

(βTi)

τ1

SiC

TiSi2

TiSi

Ti5Si

4

T2

TiCx

80 60 40 20

1000

1250

1500

1750

Ti 95.00

C 5.00

Si 0.00

Ti 0.00

C 5.00

Si 95.00Ti, at.%

Te

mp

era

ture

, °C 1473

1330

1483

1570

1485

1095

1354

Ti3Si+TiCx

(βTi)+TiCx+T2

(βTi)+TiCx

TiCx+T2

L+TiC T2

T2+Ti5Si4

L+T2 L+τ1

L+SiC

(Si)+SiC+TiSi2

L+SiC+(Si)

SiC+TiSi2+τ1

L+SiC+TiSi2

τ1+T2+TiSi2

TiSi+T2+Ti5Si4 TiSi+TiSi2+T2

Fig. 8: C-Si-Ti.


1100°C [2000Du]

Fig. 9: C-Si-Ti.

Isopleth at 5 at.% C

[2000Du]

148


MSIT®

C–Si–Ti

20 40 60 80

1000

1100

1200

1300

1400

1500

1600

1700

1800

Ti 90.00

C 10.00

Si 0.00

Ti 0.00

C 10.00

Si 90.00Si, at.%

Te

mp

era

ture

, °C

1330

14731483

1354

1095

L+TiCx

(βTi)+TiCx+T2

Ti3Si+TiCx

(βTi)+TiCx

TiCx+T2

L+τ1+T2

L+SiC

τ1+T2+TiSi2

SiC+τ1

+TiSi2(Si)+SiC+TiSi2

L+SiC+TiSi2

L+SiC+(Si)

L+T2 L+τ1

TiSi+T2+TiSi2

1000

1100

1200

1300

1400

1500

1600

1700

1800

Ti 57.00

C 10.00

Si 33.00

Ti 52.00

C 10.00

Si 38.00Si, at.%

Te

mp

era

ture

, °C

T2

L+T2 L+τ1+T2

TiCx+T2

T2+Ti5Si4

1570

1485 1483

τ1+T2+TiSi2TiSi+T2+TiSi2

T2+TiSi2TiSi+T2

TiSi+T2+Ti5Si4

3534 36 37

Fig. 10a: C-Si-Ti.


[2000Du]

Fig. 10b: C-Si-Ti.


from 33 to 38 at.% Si

[2000Du]

149


MSIT®

C–Si–Ti

20 40 60

1000

1100

1200

1300

1400

1500

1600

1700

1800

Ti 85.00

C 15.00

Si 0.00

Ti 5.00

C 15.00

Si 80.00Si, at.%

Te

mp

era

ture

, °C

L+τ1

L+SiC

(Si)+SiC+TiSi2

1330

1473

SiC+τ1+TiSi2

L+τ1+T2

τ1+T2

τ1+T2+TiSi2

TiCx+T2

L+TiCx

1354

(βTi)+TiCx

(βTi)+TiCx+T2

Ti3Si+TiCx

1483

L+SiC+(Si)

L+SiC+TiSi2

1095

1100 16001500140013001200

0

25

20

15

10

5

Temperature, °C

Pre

ssu

re,

kb

ar

stable

decomposes toTiC +Ti Si +Ti(Si,C)

x 5 3

Fig. 11: C-Si-Ti.


[2000Du]

Fig. 12: C-Si-Ti.

The region of

temperature-pressure

stability range of the

Ti3SiC0.4 phase

[1993Sam]

150


MSIT®

C–Si–Ti

743

746

745

744

514

517

516

515

0.0 1.00.80.60.40.2

a-L

att

ice

Pa

ram

ete

r,p

mc-L

att

ice

Pa

ram

ete

r,p

m

Carbon Content, x

[2000Wil1][2000Tho]

1600

1500

1400

1300

0.0 0.2 0.4 0.6 1.00.8

De

po

sitio

nTe

mp

era

ture

,°C

SiCl /(SiCl +TiCl )4 4 4

TiCx

SiC+TiCx

TiC +x

τ1

SiC+TiC +x

τ1

SiC+τ1

TiSi +2 1τ

SiC+TiSi2

SiC

τ1

Fig. 13: C-Si-Ti.

Effect of the

interstitial atom C on


Ti5Si3Cx (0< x<1)

Fig. 14: C-Si-Ti.

Relationship between

the deposited phase

and the deposition

condition [1987Got]

151


MSIT®

C–Si–Ti

0.70 0.75 0.80 0.85 0.90 0.95 1.00.55

0

180

140

100

60

20

Pre

ssu

re,

GP

a

V/V0

Fig. 15: C-Si-Ti.

Reduced volume

(V/V0) of Ti3SiC2 vs

pressure [2003Jor]

152


MSIT®

Ce–Mg–Y

Cerium – Magnesium – Yttrium

Hans Flandorfer

Literature Data

The publication of [1997Fla] is concerned with the isothermal section of the ternary Ce-Mg-Y system at

500°C including also a reinvestigation of the binary Ce-Y system. The experimental results are based on a

simultaneous critical assessment of all thermodynamic phase equilibria, crystallographic and solid solution

data obtained for the Ce-Mg-Y ternary. Earlier investigations essentially focused on two partial isotherms

at 300 and 500°C in the Mg-rich corner at more than 75 mass% Mg [1981Dri]. Three isopleths were

constructed for the very Mg-rich region: Mg-30 mass% Y to Mg - 23 mass% Ce [1981Pad], Mg - 19 mass%

Y to Mg - 16 mass% Ce [1981Dri] and Mg - 15 mass% Ce to Mg - 16 mass% Y [1981Dri]. The potential

of Ce-Mg-Y alloys with respect to intermediate valence effects on the Ce atoms was investigated in a few

selected alloys prepared from the solid solutions YxCe1-xMg3 [1985Gal] and Ce1-xYxMg [1985Pie]. Phase

equilibria in the Mg-rich region of the Ce-Mg-Y diagram (up to 30 mass% Y and up to 23 mass% Ce) have

been studied by [1981Dri] and [1981Pad] on the basis of wet chemical analysis, X-ray diffraction on powder

and flat specimens, optical microscopy and scanning electron microscopy (SEM) on flat mechanically

polished alloy surfaces. Thermal effects have been recorded using chromel-alumel thermocouples or by

differential thermal analysis (DTA; accuracy 2°C) under a protective flux of 80% LiCl+20 % LiF. All

alloys were prepared from elemental ingots with a purity better than 99.95 mass% (Mg,Y) or 99.85 mass%

(Ce) by melting in alumina crucibles in an electric resistance furnace. The rare earths were introduced via

pre-alloyed master alloys containing 30 to 40 mass% of the rare earth metal. The final melts were cast into

copper molds. After a 50 % deformation of the ingots at 480°C in steel mandrels, the specimens were

annealed in vacuum sealed quartz ampules for 50 h at 500°C and for 500 h at 300°C, respectively, prior to

quenching into water. The microstructure was also studied on slowly cooled specimens. A 30 % solution of

H3PO4 in alcohol was used as etchant.

The metals used by [1997Fla] were magnesium rod (99.9 mass%), cerium rod (99.9 mass%) and yttrium

ingots (99.9 mass%). Ingots of the pure elements were enclosed in small cylindrical tantalum crucibles

sealed by arc welding under pure argon. The samples were melted and remelted in an induction furnace

under continuous shaking of the crucibles in a stream of high purity argon. The tantalum crucibles were then

sealed in iron cylinders wherein the samples were annealed in a resistance furnace at 500°C for 72 to 168 h

and finally quenched into cold water. Thermodynamic equilibrium in the Ce-Y binary as well as in the

ternary region at low magnesium contents was only attained after prolonged heat treatments of more than

3000 h at 500°C. Some of the Ce-Y binary alloys were simply fused in an argon arc furnace prior to

annealing in silica tubes which were vacuum sealed and internally lined with protective molybdenum foil.

Specific alloys have been investigated by DTA with heating and cooling rates of 2 to 10 K min-1 with

a 1 % overall accuracy of the temperatures recorded. Light optical microscopy, scanning electron

microscopy, microprobe and X-ray analyses were used to examine phase equilibria and equilibrium

compositions [1997Fla]. The critical assessment was published in [1997Fla, 1996Fla1].

Binary Systems

The binary system Ce-Mg has been accepted from [Mas2]. The preliminary version of the Ce-Y system

presented in [Mas2] based on a detailed assessment by [1986Gsc] was revised by [1997Fla] (see Fig. 1).

According to the observations of [1997Fla] the phase ( Sm type) which was reported [Mas2] at about 50

at.% Y seems to be a metastable phase as it was found to disappear after long-term annealing for 3000 h at

500°C. The Mg-Y system has been recently revised by [1994Gio]. The work of [1995Gio] revealed an

extended homogeneity region for Y1-yMg2.

Solid Phases

153


MSIT®

Ce–Mg–Y

Phase equilibria within the entire Ce-Mg-Y system refer to the isothermal section derived at 500°C

comprising two hitherto unknown ternary compounds YxCe1-xMg5-y and YxCe1-xMg2 (see Fig. 2).

Crystallographic data for the Ce-Mg-Y ternary compounds YxCe1-xMg5-y (0.39 x 0.84; 0 y 0.60)

and YxCe1-xMg2 (x 0.67) are presented in Table 1 and are based on quantitative Rietveld intensity analyses

of X-ray powder diffractometer data [1996Fla2]. The consecutive formation of structurally closely related

phase regions as a function of Y/Ce substitution such as the pairs: Y1-yMg2 (MgZn2 type) - YxCe1-xMg2

(MgCu2 type) and Y5-yMg24+y ( Mn type) - YxCe1-xMg5-y (GdMg5 or defect Sm11Cd45 type), thereby

follows the concept of pseudolanthanoids [1995Gio] (i.e. the rule of Gschneidner and Valletta [1968Gsc]).

A listing of the experimentally observed unit cell dimensions in ternary multiphase alloys is given in

Table 1. Despite the practically nonisomorphous nature of the light rare earth and the heavy rare earth

magnesium compounds, mutual solid solubilities between cerium - and yttrium - magnesium phases are

naturally large. YMg and CeMg form a continuous solid solution at 500°C. A quite extensive solution range

is encountered for YxCe1-xMg3 (0 x 0.60). Solid solutions YxCe1-xMg and YxCe1-xMg3 closely obey

Vegard's rule with a linear variation of the unit cell dimensions indicating a rather ideal random substitution

Ce/Y. This is also true for the homogeneous region of the ternary phase (Ce1-xYx)Mg5-y. A large

homogeneous region is observed for binary Y1-yMg2 (0 y 0.32 at 500°C [1994Gio, 1995Gio]) which

continues into the ternary (YxCe1-x)1-yMg2 up to x = 0.18 for y = 0 and up to x = 0.13 for y = 0.32. Mg-rich

compounds Y5-yMg24+y, CeMg12 and Ce5Mg41 at 500°C show only small solid solubility in the ternary

(about 0.7 at.% Ce in Y5-yMg24+y, 0.5 at.% Y in CeMg12 and 1.5 at.% Y in Ce5Mg41). The solubility of the

rare earth components in magnesium were determined by [1981Dri] for two temperatures, 300°C and

500°C. At 300°C a maximum solubility of 1.36 at.% Y and 0.01 at.% Ce was reported [1981Dri]. At 500°C

these authors reported 2.92 at.% Y and 0.05 at.% Ce in satisfactory agreement with EPMA data obtained

by [1996Fla1, 1997Fla].


DTA experiments, supported also by micrographic analyses, revealed an incongruent melting behavior for

YxCe1-xMg5-y at (597°C at x = 0.6 and y = 0.3) as well as for YxCe1-xMg2 with the MgCu2 type at x 0.67

(760°C). The finely grained ternary eutectic in the Mg-rich region, L (Mg)+CeMg12+Y5Mg24, first

reported [1981Dri, 1981Pad] at 543 2°C has been confirmed by [1997Fla] from DTA experiments

(542 3°C) and metallographic observations. On the basis of these data, a reaction scheme was constructed

for the Mg-rich part of the system (see Fig. 3)

Liquidus Surface

No complete liquidus surface has been provided so far. However, a partial liquidus projection for the

Mg-rich corner (< 85 at.% Mg) is available from earlier investigations [1981Dri, 1981Pad] and is shown in

Fig. 4.

Isothermal Sections

Figure 2 shows the isothermal section at 500°C after [1997Fla]. Phase triangulation and solid solubility

limits referring to the Mg-rich corner of the diagram are consistent with the data by [1981Dri]. Phase

equilibria are characterized by the formation of extended or complete solid solutions and by the existence

of two ternary compounds. Determination of the vertices of the three phase equilibria in the Mg-poor region

by EMPA is severely hampered by the extremely fine grained microstructure and/or finely dispersed

substructure of the grains. Therefore, this part of the diagram (<50 at.% Mg) was essentially based on the data

of [1997Fla] and is presented by dashed lines (Fig. 2).

Miscellaneous

Three isopleths have been established in the Mg-rich region of the ternary system confirming the

coexistence of the Mg solid solution with CeMg12 and Y5-yMg24+y: at 19 mass% Y from 0 to 16 mass% Ce

154


MSIT®

Ce–Mg–Y

intersecting the ternary eutectic composition (Fig. 5, after [1981Dri]), isopleth at 15 mass% Ce for 0 to 20

mass% Y (Fig. 6, after [1981Dri]), and the section 30Y70Mg-23Ce77Mg (Fig. 7 after [1981Pad]).

Magnetic and electrical behavior of the binary Ce-Mg compounds, as well as of the solid solutions

(Ce,Y)Mg and (Ce,Y)Mg3 and the ternary phases (Ce,Y)Mg2 and (Ce,Y)Mg5 were reported by [1996Fla2].

References

[1965Woo] Wood, D.H., Cramer, E.M., “Phase Relations in the Magnesium-Rich Portion of the

Cerium-Magnesium System”, J. Less-Common Met., 9, 321-337 (1965) (Equi. Diagram,

Crys. Structure, Experimental, 13)

[1968Gsc] Gschneidner, K.A., Valletta, R.M., “Concerning the Crystal Structure Sequence in the

Lanthanide Metals and Alloys; Evidence for 4f Contribution to the Bonding”, Acta. Met.,

16, 477, (1968) (Theory, 47)

[1979Dar] Darriet, B., Pezat, M., Hbika, A., “Magnesium-Rich Alloys of the Rare Earth and their

Application to Hydrogen Storage” (in French), Mater. Res. Bull., 14, 377-385 (1979) (Crys.


[1981Dri] Drits, M.E., Padezhnova, E.M., Dobatkina, T.V., Votekhova, E.A., Kinzhibalo, V.V., “The

Magnesium Corner of the Mg-Y-Ce System”, Russ. Metall., 6, 200-203 (1981), translated

from Izv. Akad. Nauk SSSR, Met., 6, 206-210 (1981) (Equi. Diagram, Crys. Structure,

Experimental, #, 3)

[1981Pad] Padezhnova, E.M., Melnik, E.V., Kinzhibalo, V.V., Dobatkina, T.V., “Nature of the

Interaction of Elements in Magnesium Alloys of the Mg-Y-Ce System”, Russ. Metall., 6,

203-206 (1981), translated from Izv. Akad. Nauk SSSR, Met., 4, 220-223 (1981) (Equi.

Diagram, Crys. Structure, Experimental, #, 7)

[1985Gal] Galera, R.M., Pierre, J., Murani, A.P., “Spin Dynamics in CexR1-xMg and CexR1 xMg3

(R = La, Y)”, J. Magn. Magn. Mater., 47-48, 155-158 (1985) (Experimental, 12)

[1985Pie] Pierre, J., Galera, R.M., Siaud, E., “Evidence for Kondo-Type Behaviour in CexR1-xM

Compounds with R = La, Y and M = Mg, Zn”, J. Phys. (Paris), 46, 621-626 (1985)

(Experimental, 13)

[1986Gsc] Gschneidner, K.A., Calderwood, F.W., “Intra Rare Earth Binary Alloys: Phase

Relationship, Lattice Parameters and Systematics” in “Handbook on the Physics and

Chemistry of Rare Earths”, 8, Gschneidner, Jr.K.A., Eyring, L.R., (Eds.), North Holland,

Amsterdam, 110-118 (1986) (Review, 10)

[1986Nay] Nayeb-Hashemi, A.A., Clark, J.B., “The Ce-Mg System”, Bull. Alloy Phase Diagrams,

9(2), 162-171 (1986) (Equi. Diagram, Review, 58)

[1994Gio] Giovannini, M., Marazza, R., Saccone, A., Ferro, R., “Isothermal Section from 50 to 75 at.%

Mg of the Ternary System Y-La-Mg”, J. Alloys Compd., 203, 177-180 (1994) (Equi.

Diagram, Crys. Structure, Experimental, Review, 31)

[1995Gio] Giovannini, M., Marazza, R., Saccone, A., Ferro, R., “The Isothermal Section at 500°C of

the Y-La-Mg Ternary System”, Met. Trans., 26A, 5-10 (1995) (Equi. Diagram, Crys.

Structure, Experimental, Review, 28)

[1996Fla1] Flandorfer, H., “Cerium-Magnesium-Yttrium”, in “Final Report MSIT Training Network on

Constitution of Engineering Materials”, Effenberg, G., (Ed.), Contract: CHRX -

CT93-0285, European Community Program TMR; Human Capital and Mobility, (1996)


[1996Fla2] Flandorfer, H., Kostikas, A., Godart, C., Giovannini, M., Saccone, A., Ferro, R., “On the

Magnetic and Valence Properties of Ce-Mg-Y Compounds”, J. Alloys Compd., 240,


[1997Fla] Flandorfer, H., Giovannini, M., Saccone, A., Rogl, P., Ferro, R., “The Ce-Mg-Y System”,

Metall. Trans. A, 28(2), 265 276 (1997) (Equi. Diagram, Experimental, Review, *, 24)

155


MSIT®

Ce–Mg–Y


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Ce(

< 798

cI2

Im3m

W

a = 412 [Mas2]

( Ce)

< 726

cF4

Fm3m

Cu

a = 516.1 [Mas2]

( Ce)

< 61

hP4

P63/mmc

La

a = 368.10

c = 1185.7

[Mas2]

( Ce)

< 177

cF4

Cu

a = 485 [Mas2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.076

a = 323.5

c = 522.0

[Mas2]

at 3.6 at.% Y,

0.1 at.% Ce [1997Fla]

, (Y)

< 1522

cI2

Im3m

W

a = 407 [Mas2]

, (Y)

< 1478

hP2

P63/mmc

Mg

a = 364.82

c = 573.18

[Mas2]

(Ce,Y) hR3

R3m

Sm

a = 365.88

c = 2638.4

a = 366.03

c = 2639.5

at 54.17 at.% Ce

at 55.89 at.% Ce

metastable phase [1997Fla]

YxCe1-xMg

CeMg

< 711

Y1-yMg

< 935

cP2

Pm3m

CsCl

cP2

Pm3m

CsCl

a = 389.1

a = 381.9

a = 390.8

a = 381

a = 378.1

0 x 1

at x = 0.2

at x = 0.8 [1997Fla]

[Mas2]

[V-C2]

[Mas2]

at 50 at.% Mg

at 52 at.% Mg [V-C]

* 1, YxCe1-xMg2

< 760

CeMg2

750 - 615

cF24

Fd3m

Cu2Mg

a = 860.7

a = 873.3

at x = 0.67

[1997Fla]

[Mas2]

[1986Nay]

156


MSIT®

Ce–Mg–Y

(Y1-xCex)1-yMg2

Y1-yMg2

< 780

hP12

P63/mmc

MgZn2

a = 604.7

c = 980.8

a = 603.8

c = 979.0

a = 603.4

c = 978.6

a = 603.0

c = 977.2

a = 602.7

c = 976.3

a = 602.3

c = 974.7

at x = 0.11, y = 0.23

0 x 0.18 for y = 0

0 x 0.13 for y = 0.32 [1997Fla]

0 y 0.3 [Mas2]

at 66.7 at.% Mg

[1994Gio]

at 67.95 at.% Mg

[1994Gio]

at 71 99 at.% Mg

[1994Gio]

at 72.97 at.% Mg

[1994Gio]

at 73.82 at.% Mg

[1994Gio]

YxCe1-xMg3

CeMg3

< 798

cF16

Fm3m

BiF3

a = 742.7

a = 741.3

a = 744.3

0 x 0.6

at x = 0.10

at x = 0.30 [1997Fla]

[Mas2]

[V-C2]

* 2, (YxCe1-x)Mg5-y

< 597

cF448

F43m

Sm11Cd45 a = 2248.3

a = 2245.8

a = 2244.0

a = 2243.2

a = 2238.8

0.39 x 0.89

0 y 0.6 [1997Fla]

at x = 0.4, y = 0.3

at x = 0.51, y = 0.3

at x = 0.60, y = 0.3

at x = 0.63, y = 0.3

at x = 0.77, y = 0.3

Ce5Mg41

< 635

tI92

I4/m

Ce5Mg41

a = 1478

c = 1043

[Mas2]

[V-C2]

CeMg10.3

621 - 611

hP38

P63/mmc

Th2Ni17

a = 1033

c = 1025

[1986Nay]

CeMg12 ( )

< 616

tI26

I4/mmm

ThMn12

a = 1033

c = 596

[Mas2]

[V-C2]

CeMg12 ( ) oI338

Immm

CeMg12 ( )

a = b = 1033

c = 7750

a = b = 1032.1

c = 7701

a = b = 1032.4

c = 7723

a = b = 1033.7

c = 7737

a = b = 1032.0

c = 7724

at 7.69 at.% Ce

[1965Woo]

at 8.85 at.% Ce

[1979Dar]

at 8.33 at.% Ce

[1979Dar]

at 7.69 at.% Ce

[1979Dar]

at 7.14 at.% Ce

[1979Dar]

Y5-yMg24+y

< 605

cI58

I43m

Mn

a = 1127.8

a = 1125.0

0 y 1.23 [Mas2]

at 84 at.% Mg

at 87 at.% Mg [V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

157


MSIT®

Ce–Mg–Y



Ce Y Mg

L Mg+CeMg12+Y5-yMg24+y 542

[1996Fla2]

543

[1981Dri]

E L

CeMg12

Y5-yMg24+y

(Mg)

2.5

0.5

0.5

0.06

6.7

7.5

12.5

3.6

90.8

92

87

96.3

20 40 60 80

0

250

500

750

1000

1250

1500

Y Ce

Ce, at.%

Te

mp

era

ture

, °C

L

1478°C

~ 780°C

726°C~ 705°C

1527°C

798°C

L + (δCe, βY)

(δCe, βY)

(δCe, βY) + (αY)

(αY) + (βCe)

(γCe)

(αY)

(βCe)

(βCe) + (γCe)

Fig. 1: Ce-Mg-Y.

Binary system Ce-Y

158


MSIT®

Ce–Mg–Y

20

40

60

80

20 40 60 80

20

40

60

80

Y Ce

Mg Data / Grid: at.%

Axes: at.%

(γCe)(αY)

YMg

(βCe)

CeMg

Ce5Mg

41

CeMg3

τ1

CeMg12

τ2

Y5Mg

24

Y1-y

Mg2

YxCe

1-xMg

Fig. 2: Ce-Mg-Y.


500°C

Fig. 3: Ce-Mg-Y. Partial reaction scheme

Ce-Mg A-B-CCe-Mg-Y Mg-Y

l (Mg)+CeMg12

(β)

592 e1

L (Mg)+CeMg12

(β)+Y5Mg

24543 E

1

(Mg)+CeMg12

(β)+Y5Mg

24

l (Mg)+Y5Mg

24

566 e2L+CeMg

12(β)+Y

5Mg

24

159


MSIT®

Ce–Mg–Y

10

10

90

Y 15.00

Ce 0.00

Mg 85.00

Y 0.00

Ce 15.00

Mg 85.00


Axes: at.%

(Mg)

E1,543e

2,566

e1,592

Y5-y

Mg24+y

CeMg12

(β)

500

600

Ce 0.00

Mg 94.00

Y 6.00

Ce 3.80

Mg 89.10

Y 7.10Ce, at.%

Te

mp

era

ture

, °C

L

L+(Mg)+Y5Mg24

L + (Mg)

(Mg)+

Y5M

g2

4

(Mg) + Y5Mg24+x+ CeMg12

L+CeMg12

L+Y5Mg24+CeMg12543

E1

Fig. 4: Ce-Mg-Y.

Partial projection of

the liquidus surface in

the Mg corner

Fig. 5: Ce-Mg-Y.

Partial isopleth at 19

mass% Y, from 0 to

16 mass% Ce

160


MSIT®

Ce–Mg–Y

500

600

Ce 0.00

Mg 89.51

Y 10.49

Ce 4.93

Mg 95.07

Y 0.00Ce, at.%

Te

mp

era

ture

, °C

L

L+(Mg)+Y5Mg24

L + (Mg)

(Mg)+Y5Mg24

(Mg) + Y5Mg24+ CeMg12

L+CeMg12

L+Y5Mg24

543

L+(Mg)+CeMg12

(Mg)+CeMg12

500

600

Ce 2.97

Mg 97.03

Y 0.00

Ce 3.56

Mg 88.95

Y 7.48Y, at.%

Te

mp

era

ture

, °C

L

L + (Mg)

543L + (Mg) + CeMg12

(Mg) + CeMg12

L + CeMg12

L + Y5Mg24 + CeMg12

(Mg) + Y5Mg24+ CeMg12

592°C

630°C

Fig. 7: Ce-Mg-Y.

Section 30Y70Mg

-23Ce77Mg (in

mass%)

Fig. 6: Ce-Mg-Y.


mass% Ce, from 0 to

20 mass% Y

161


MSIT®

Ce–Mg–Zn

Cerium – Magnesium – Zinc

Uwe Kolitsch, Peter Bellen, Stefanie Kaesche, Daniele Macciò, Natalia Bochvar, Yurii Liberov, Peter Rogl

Literature Data

The system was first studied in the region between 0 and 80 mass% Mg by [1946Kor]. Starting materials

were Mg (purity 99.8 to 99.9 %), Zn of undefined purity, and Ce (purity 96 %). The alloys were fused in

alumina crucibles under a protective flux of KCl-LiCl. Ce was introduced as a Ce-Mg prealloy containing

21 to 22 % Ce. Thermal analysis was used to construct six vertical sections with a constant mass ratio Zn:Ce

of 1:5, 1:2, 1:1, 2:1, 4.5:1 and 10:1, and to draw a liquidus surface of the Mg corner. Samples were annealed

at 335°C for 240 h, at 300°C for 360 h, and at 200°C for 480 h. Some samples were allowed to cool “slowly”

[1946Kor] in the furnace. The annealed samples were water quenched, and subsequently etched with 2 %

alcoholic nitric acid. From the microstructure of the etched samples, the combined solubility of Zn and Ce

in Mg was determined.

Subsequent reports [1977Mel, 1978Mel, 1983Mel, 1989Dri] studied the phase relations in the Ce-Mg-Zn

ternary system in the region Mg-MgZn2-CeMg-CeZn. The alloys used in the experiments were prepared by

fusing the elements in crucibles ([1989Dri] states alumina crucibles) under a protective layer of VI2 flux.

The starting materials were metal ingots of purity better than 99.5 mass% [1977Mel, 1978Mel, 1989Dri].

Phase triangulation in the region Mg-MgZn2-CeMg-CeZn has been evaluated by means of X-ray powder

diffraction [1978Mel] from 150 alloys annealed at 300°C for 240 h and quenched into water. Four ternary

compounds were observed, 1 to 4, the crystal structures of 1 and 3 being unknown. [1989Dri]

constructed two polythermal sections in the Mg rich corner at 24 mass% Zn and 34 mass% Zn, both ranging

from 0 to 20 mass% Ce. Methods used were DTA (cooling rates 2 to 5 K min-1), EMPA, SEM and X-ray

powder diffraction. Microstructure was investigated after mechanical polishing, applying a solution of 30 %

H3PO4 in alcohol as etchant.

[1980Zak], in his review, plotted, based on [1946Kor], the curves of the combined solubility of Zn and Ce

in (Mg).

Binary Systems

The accepted Mg-Zn binary system has been calculated on the basis of a critical assessment [1992Aga]. The

Ce-Mg and Ce-Zn binary systems have been accepted from [Mas2]. The crystallographic data relevant to

the binary boundary systems are listed in Table 1.

Solid Phases

Four ternary compounds 1, 2, 3, and 4 have been observed, of which 2 and 4 exhibit extended solid

solution ranges at constant Ce contents [1978Mel]. The lattice parameters of the solid solution ranges are

given in Table 1. [1989Dri] assumed that 2 derives from the isostructural binary compound CeMg10.3.


[1946Kor] found a ternary eutectic at 341 to 343°C with a composition corresponding to 50 mass% Mg,

47.5 mass% Zn, and 2.5 mass% Ce. On the basis of the two experimentally established isothermal reactions

L+ 2 (Mg)+ 1 at 349 1°C and L (Mg)+ 1+Mg51Zn20 at 341 1°C [1989Dri], and the eutectoid ternary

decomposition of Mg51Zn20 mentioned by [1989Dri], a partial reaction scheme has been derived of the

region close to the binary phases MgZn and Mg51Zn20 (Fig. 1). The reaction scheme requires a further

transition reaction rather close to the ternary eutectic at 340°C. No experimental details are hitherto

available regarding the sequence of the two reactions, which is due to the extremely close region of their

appearance. Therefore, these two reactions were integrated into a degenerate reaction D1,2 at ~340°C

slightly below the binary formation temperature of Mg51Zn20.

162


MSIT®

Ce–Mg–Zn

Liquidus Surface

A projection of the liquidus surface of the region between 0 and 80 mass% Mg showing the ternary eutectic

E1 has been given by [1946Kor]. It is not reproduced here for two reasons: firstly, Ce used was of inadequate

purity. Secondly, the location of the eutectic trough starting at the Ce-Mg binary and joining E1 is only

approximate according to [1946Kor].

Isothermal Sections

[1946Kor] gave a projection of the solubility isotherms for Mg-rich samples annealed at 335, 300 and

200°C, and for samples allowed to cool “slowly” to 20°C. Because of the low purity of Ce [1946Kor] used,

the values of the maximum combined solubilities of (Ce+Zn) in Mg (e.g., 4.9 mass% at 335°C) have to be

considered only approximate. Nonetheless, it is evident that the combined solubility rapidly increases with

increasing Zn content.

A partial isothermal section at 300°C has been evaluated by [1978Mel] for the region

Mg-MgZn2-CeMg-CeZn (see Fig. 2). The tie line between 1 and Mg51Zn20 given by [1978Mel] has been

omitted because Mg51Zn20 is not stable at 300°C [1992Aga]. Phase equilibria in this section are

characterized by the existence of extended binary and ternary solid solutions, each at a constant Ce content.


Two vertical sections from Mg to 80 mass% Ce10Zn and from Mg to 80 mass% CeZn are given by

[1946Kor]. Although they were found to approximately comply with later data, they are not reproduced here

due to the low purity Ce [1946Kor] used. Two polythermal sections have also been established by [1989Dri]

at 24 mass% Zn and 34 mass% Zn, both ranging from 0 to 20 mass% Ce. Both sections are reproduced in

Figs. 3 and 4, respectively, with small changes to comply with the exact location of the Ce-poor boundary

of the (Mg)+ 2 two phase field in the isothermal section at 300°C established by [1978Mel]. Details

concerning the area near to the Mg-Zn binary system are schematically shown in Fig. 5.

Miscellaneous

Ageing, hardness, and corrosion behavior of alloys were studied by [1946Kor].

Further investigations in this system should focus on structure determination of the ternary compounds 1

and 3, and on experimental investigation of the unknown regions of the phase diagram.

References

[1946Kor] Korolkov, A.M., Saldau, Ya.P., “Solubility of Zn and Ce in Mg in the Solid State” (in

Russian), Izv. Sekt. Fiz.-Khim. Anal., 16(2), 295-306 (1946) (Equi. Diagram,

Experimental, 18)

[1959Ray] Raynor, G.V., “Intermediate Phases in Magnesium Alloys” in “The Physical Metallurgy of

Magnesium and Its Alloys”, Pergamon Press, London- New York- Paris- Los Angeles,

Chapter 6, 145-215 (1959) (Review, Equi. Diagram, Crys. Structure, 35)

[1965Woo] Wood, D.H., Kramer, E.M., “Phase Relations in the Magnesium-Rich Portion of the

Cerium-Magnesium System”, J. Less-Common Met., 9, 321-337 (1965) (Equi. Diagram,

Crys. Structure, Experimental)

[1977Mel] Melnik, E.V., Zmii, O.F., Cherkasim, E.E., “On the Structure of the Ce2Mg3Zn3

Compound” (in Russian), Vestn. L`viv. Univ. Ser. Khim., 19, 34-36 (1977) (Crys.

Structure, 5)

[1978Mel] Melnik, E.V., Kostina, M.F., Yarmlyuk, Ya.P., Zmii, O.F., “Study of the

Magnesium-Zinc-Cerium and Magnesium-Zinc-Calcium Ternary Systems” (in Russian),

Magnievye Splavy, Mater.Vses. Soveshch. Issled., Razrab. Primen. Magnievyhk Splavov,

95-99 (1978) (Equi. Diagram, Crys. Structure, Experimental, #, *, 15)

163


MSIT®

Ce–Mg–Zn

[1979Dar] Darriet, B., Pezat, M., Hbika, A., Hagenmuller, P., “Rare Earth Compounds with

Magnesium and Their Application in Hydrogen Storage” (in French), Mater. Res. Bull.,

14(3), 377-385 (1979) (Equi. Diagram, Crys. Structure, Experimental)

[1980Zak] Zakharov, A.M., “High-Strength Alloys of the Mg-Zn-Zr System” (in Russian), in

“Promyschlennye Splavy Tsvetnykh Metallov”, Moscow, Metallurgiya, 101-104 (1980)

(Equi. Diagram, Review, *, 4)

[1983Mel] Mel'nik, E.V., Zmiy, O.F., Muratova, E.B., “Interaction Between IMC in Ternary

Mg-Zn-(In, La, Ce) Systems” (in Russian), IV Vsesoyuzn. Konfer. Kristallokhimii

Intermetallich. Soyedin., Tezisy Doklad., L'vov Univ., L'vov, 38 (1983) (Crys. Structure,

Experimental, 0)

[1986Kin] Kinzhibalo, V.V., Tyvanchuk, A.T., Mel'nik, E.V., “Crystal Structure of the Compounds of

Magnesium and Zinc with Rare-Earth Metals and Calcium”, IV Vsesoyuzn. Soveshch. po

Kristallokhimii Neorgan. i Koordin. Soyedin., Tezisy Doklad., Moscow, Nauka, 196 (1986)


[1989Dri] Drits, M.E., Drozdova, E.I., Korolkova, I.G., Kinzhibalo, V.V., Tyvanchuk, A.T.,

“Investigation of Polythermal Sections of the Mg-Zn-Ce System in the Magnesium-Rich

Region”, Russ. Metall., 2, 195-197 (1989), translated from Izv. Akad. Nauk SSSR, Met., 2,

198-200 (1989) (Equi. Diagram, Experimental, Crys. Structure, #, *, 9)

[1992Aga] Agarwal, R., Fries, S.G., Lukas, H.L., Petzow, G., Sommer, F., Chart, T.G., Effenberg, G.,

“Assessment of the Mg-Zn System”, Z. Metallkd., 83(4), 216-223 (1992) (Equi. Diagram,

Thermodyn., Review, #, 44)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

a = 320.99

c = 521.08

a = 319.57

c = 518.82

[Mas2]

[V-C2]

at 2.8 at.% Zn [V-C2]

linear dependency

Ce5Mg41

< 635

tI92

I4/m

Ce5Mg41

a = 1478

c = 1043

[Mas2, V-C2]

CeMg10.3

621-611

hP38

P63/mmc

Th2Ni17

a = 1033

c = 1025

[Mas2, V-C2]

164


MSIT®

Ce–Mg–Zn

CeMg12 ( )

< 616

oI338

Immm

CeMg12 ( ) a = 1033

b = 1033

c = 7750

a = 1032.1

b = 1032.1

c = 7701

a = 1032.4

b = 1032.4

c = 7723

a = 1033.7

b = 1033.7

c = 7737

a = 1032.0

b = 1032.0

c = 7724

Pseudotetragonal long-range order

variant of CeMg12 [1965Woo]

at 7.69 at.% Ce [1965Woo]

at 8.85 at.% Ce [1979Dar]

at 8.33 at.% Ce [1979Dar]

at 7.69 at.% Ce [1979Dar]

at 7.14 at.% Ce [1979Dar]

Mg51Zn20

325-341

oI158

Immm

Mg51Zn20

a = 1408.3

b = 1448.6

c = 1402.5

[V-C2]

Composition Mg0.718Zn0.282 [1992Aga];

labelled “Mg7Zn3” in [Mas2]

MgZn

< 347

oP48

?

MgZn

a = 923

b = 533

c = 1760

Composition Mg0.48Zn0.52 [1992Aga]

[Mas2, 1959Ray]

Mg2Zn3

< 416

mC110

C2/m

Mg4Zn7

a = 2596

b = 524

c = 2678

= 148.6°

[V-C2],

labelled Mg4Zn7

Composition Mg0.40Zn0.60 [1992Aga]

MgZn2

< 587

hP12

P63/mmc

MgZn2

a = 522.3

c = 856.6

[1992Aga, V-C2]

CeMg1-xZnx

CeMg

< 711

CeZn

< 825

cP2

Pm3m

CsCl

a = 390.8

c = 390

a = 370.4

c = 371.3

0 x 1 at 300°C [1978Mel]

linear dependency, scaled from diagram

[Mas2]

[V-C2], at 25°C

[1978Mel], quenched from 300°C

[Mas2]

[V-C2]

[1978Mel]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

165


MSIT®

Ce–Mg–Zn

Ce(Mg1-xZnx)12 ( )

CeMg12 ( )

< 616

tI26

I4/mmm

ThMn12

a = 1033

c = 596

0 x 0.08 at 300°C [1978Mel]

[Mas2] [V-C2]

Ce(Mg1-xZnx)3

CeMg3

< 798

cF16

Fm3m

BiF3 a = 744.3

a = 742.8

a = 707.6

[Mas2]; 0 x 0.39 at 300°C

[1978Mel], linear dependence

at x = 0, at 300°C [V-C2]

at x = 0, quenched from 300°C

[1978Mel]

at x = 0.39, quenched from 300°C

[1978Mel]

* 1, CeMg7Zn12 h**

?

a = 1471.0

c = 880.0

[1978Mel, 1989Dri, 1986Kin]

* 2, Ce(Mg,Zn)10.1 hP38

P63/mmc

Th2Ni17

a = 1010

c = 997

a = 960

c = 947

from 9.1 to 45.5 at.% Zn at 300°C;

possibly related to CeMg10.3 [1989Dri]

labelled Ce(Mg,Zn)9 by [1978Mel]

at Ce(Mg0.9Zn0.1)10.1 [1989Dri]

at Ce(Mg0.5Zn0.5)10.1 [1989Dri]

* 3, CeMg3Zn5 ? ? [1978Mel]

* 4, Ce(Mg,Zn)3 cF16

Li2AgMg or

MnCu2Al a = 706.4 0.4

a = 701.1

a = 708.9

from 35 to 45 at.% Zn at 300°C

[1978Mel]

for CeMg1.5Zn1.5 [1977Mel, 1978Mel]

for CeMg1.2Zn1.8 [1978Mel]

for CeMg1.6Zn1.4 [1978Mel]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

166


MSIT®

Ce–Mg–Zn

Mg-Zn Ce-Mg-Zn

l+(Mg) Mg51

Zn20

341.0 p1

l MgZn+Mg51

Zn20

340.9 e1

Mg51

Zn20

(Mg)+MgZn

325 e2

L+τ2

(Mg)+τ1

349 U1

L, (Mg), MgZn, Mg51

Zn20

, τ1

~340 D1,2

Mg51

Zn20

τ1+(Mg)+MgZn325 E

1

L+(Mg)+τ1

L+(Mg)+τ2

τ2+L+τ

1

τ1+(Mg)+MgZn

τ2+(Mg)+τ

1

L+MgZn+τ1

MgZn+Mg51

Zn20

+τ1

(Mg)+Mg51

Zn20

+τ1

20

40

60

80

20 40 60 80

20

40

60

80

Ce Mg

Zn Data / Grid: at.%

Axes: at.%

MgZn2

τ3

τ1

Mg2Zn

3

MgZn

τ4

τ2

CeMgCeMg

12Ce5Mg

41CeMg

3

CeZn

(Mg)

Fig. 2: Ce-Mg-Zn.

Partial isothermal

section at 300°C

Fig. 1: Ce-Mg-Zn.

Reaction scheme

167


MSIT®

Ce–Mg–Zn

300

400

500

600

Ce 0.00

Mg 89.49

Zn 10.51

Ce 5.07

Mg 81.88

Zn 13.05Ce, at.%

Te

mp

era

ture

, °C

L

L + (Mg)

(Mg)+Mg51Zn20+τ1

(Mg)+τ1+τ2

(Mg)+MgZn+τ1

349°C

< 325°C

L+Mg+τ2

L+(Mg)+τ1

(Mg)+τ2

Fig. 3: Ce-Mg-Zn.

Polythermal section at

24 mass% Zn

300

400

500

600

Ce 0.00

Mg 83.93

Zn 16.07

Ce 3.93

Mg 76.99

Zn 19.08Ce, at.%

Te

mp

era

ture

, °C

L

L + (Mg)

L + (Mg) + τ2

L + (Mg) +τ1

(Mg)+Mg51Zn20+τ1 (Mg) + τ1+ τ2

(Mg)+MgZn+τ1

349°C

< 325°C

Fig. 4: Ce-Mg-Zn.


34 mass% Zn

168


MSIT®

Ce–Mg–Zn

300

400

Ce 0.00

Mg 83.93

Zn 16.07

Ce 1.16

Mg 81.87

Zn 16.96Ce, at.%

Te

mp

era

ture

, °C L + (Mg)

L + (Mg) + τ2

L + (Mg) +τ1

(Mg)+Mg51Zn20+τ1

(Mg)+MgZn+τ1

349°C

< 325°C

~ 340°C

(Mg) + MgZn

(Mg) + MgZn + Mg51Zn20

(Mg) + Mg51Zn20

341°C

325°C

0.5 1.0

Fig. 5: Ce-Mg-Zn.

Schematic

polythermal section at

34 mass% Zn near

Mg-Zn side

169


MSIT®

Co–Ni–Ti

Cobalt – Nickel – Titanium

Gabriele Cacciamani and Paola Riani

Literature Data

The Co-Ni-Ti system is characterized by the presence of one ternary and several binary Co-Ti and Ni-Ti

phases with more or less extended ternary solubility. Phase equilibria have been studied by different authors

[1975Kor, 1979Sag, 1980Gry, 1981Loo, 1983Gry, 1986Liu, 1993Du] who proposed results sometime

contradictory. One of the investigations here considered more reliable was carried out by [1981Loo]. They

determined the 900°C isothermal section by preparing 14 diffusion couples and 20 binary and ternary

samples starting from 99.97 (for Ti) and 99.99 mass% (for Co and Ni) pure elements. Alloys were arc

melted under argon and homogenized for at least 170 h at 900°C. Diffusion couples were hot pressed for at

least 64 h in a vacuum better than 10-3 Pa. Couples and alloys were then investigated by optical and electron

microscopy, microprobe analysis and X-ray diffraction.

An assessment of this system including 12 references has been recently compiled by [1999Gup].

Binary Systems

The Co-Ni system is accepted from the assessment by [1991Nis]. The Co-Ti phase diagram and structure

data are accepted from the thermodynamic assessment by [2000Cac] except for the TiCo melting

temperature, more recently determined by [2001Dav]. [1999Zha] proposed a tentative phase diagram at low

temperature involving the ( Co), ( Co) and TiCo3 phases. Due to scattered experimental data, the

extrapolated phase diagram was not retained in this assessment. The binary Ni-Ti system has been

extensively reviewed by [1991Mur]. More recently, [1996Bel] has done a new assessment of the

thermodynamic properties of the stable phases, based on thermochemical and phase diagram data from the

literature. Their calculation is in good agreement with the phase equilibria reported by [1991Mur]. The solid

homogeneity range of TiNi3 has been reproduced by [1996Bel] and is in good agreement with the literature

data. The accepted diagram in this assessment is then a compilation of [1991Mur] and [1996Bel]. Crystal

structure data and metastable phases and martensitic transformations are described by [1991Mur].

Solid Phases

Binary Co-Ti and Ni-Ti phases generally show extended ternary solubility due to the mutual substitution of

Co and Ni. In particular, for Ti(Co,Ni) and Ti2(Co,Ni), the solubility is complete [1980Gry, 1981Loo,

1983Gry, 1993Du] since the corresponding binary compounds have the same structure data. For the other

binary compounds different solubility ranges have been reported by different authors. Especially in the case

of [1981Loo] and [1993Du] differences cannot be justified by the temperature difference between the two

investigations (900 and 850°C, respectively). As an example, the solubility lobe of Ni in TiCo3 was found

to be 15 at.% Ni at 1050°C [1986Liu]. Its becomes 22 at.% Ni at 900°C [1981Loo] and 6 at.% Ni at 850°C

[1993Du].

The ternary phase Ti(Co0.5Ni0.5)3, reported by [1966Vuc, 1969Sin, 1981Loo] was not detected by

[1993Du], which however did not investigate in detail the composition range where it is stable. [1980Gry]

and [1983Gry] did not find ternary compounds. It has to be noted, moreover, that the crystal structure of the

ternary phase is intermediate between the structures of the binary compounds with the same Ti content

(TiNi3, VCo3 and AuCu3 type structures can be described in terms of different staking sequences of the

same hexagonal planes) and generates diffraction patterns which can be easily misinterpreted as due to a

superposition of the two binary phases.

[1988Kra] detected a second order phase transition which has been attributed by the authors to a martensitic

transformation in an alloy of nominal composition TiCo0.15Ni0.85. The martensitic transformations

occurring in the Ni-rich side of Ti(Co,Ni) and the formation of the R phase in the same composition area

have been investigated by [1992Shi, 1992Gou, 1999Lek].

170


MSIT®

Co–Ni–Ti

The martensitic transformation between Co and Co and the Curie temperature in Co have been studied

by [1985Shi] in the 0-25 at.% Ni, 0-10 at.% Ti composition range.

Crystal structure data for the stable and metastable Co-Ni-Ti phases are summarized in Table 1.


The quasibinary section at 50 at.% Ti is reported in Fig. 1. It is mainly based on [1975Kor] with minor

modifications in order to be consistent with the accepted congruent melting temperatures of the end

members.

Liquidus Surface

From the binary systems, the isothermal section determined by [1980Gry, 1993Du, 1981Loo, 1983Gry] and

the pseudobinary system of [1975Kor], [1999Gup] gave a speculated liquidus surface which has not be

retained in this assessment since no experimental data are reported in the literature.

Isothermal Sections

Four isothermal sections have been investigated: at 800 [1979Sag, 1980Gry], 850 [1993Du], 900 [1981Loo]

and 1000°C [1983Gry]. All of them present consistent equilibria in the Ti-richer part (at x(Ti) > 0.5). At

lower Ti content errors and contradictory results appear, especially concerning the phases at the Ti(Co,Ni)3

ratio.

At about 25 at.% Ti [1979Sag, 1980Gry] and [1983Gry] reported a continuous solid solution between TiNi3and TiCo3, which is clearly inconsistent considering the different structure of the two phases, D024 and L12

respectively. The non complete mutual solubility between these two compounds has been confirmed by

[1993Du, 1981Loo, 1986Liu]. [1979Sag, 1980Gry, 1983Gry] also indicated only one phase at the TiCo2

composition, with no ternary solubility.

The investigation by [1993Du] was not detailed in the area at < 50 at.% Ti and intermediate Co and Ni

concentrations. This is probably the reason why the ternary phase was not detected. They also suggested

solubility ranges and equilibria between the TiNi3, TiCo3 TiCo2(hexagonal) and TiCo2(cubic) phases

incompatible with the results by [1981Loo]. The section at 900°C by [1981Loo] seems to be the most

reliable, considering the accuracy of the investigation (total number of investigated samples, preparation

and analysis of both fixed composition and diffusion couple samples, etc.). In particular, at about 25 at.%

Ti, a sequence of three phases is reported with quite extended solubility ranges and narrow two-phase fields

between them (this could also explain the erroneous indication of a single solid solution reported by

[1979Sag, 1980Gry] and [1983Gry]). The isothermal section reported in Fig. 2 is then mainly bases on

[1981Loo].

Thermodynamics

The specific heat of the Ti0.5(Co1-xNix) and the Ti(Co1-yNiy) alloys has been measured by [1986Kra], at

x = 0, 0.15, 0.25, 0.5, 0.67, 1.0, and by [1987Kra, 1988Kra] at y = 0.15, in the 78-273 K temperature range.

From heat capacity measurements, the entropy was extrapolated by [1987Kra] using thermodynamic

calculations. Small values of Sfexc were found suggesting that the entropies of formation of the ternary

intermetallics are close to those for the ideal configuration. This is in agreement with the relatively small

value of Hf [1986Zas] suggesting that the solutions of Ni0.5Ti0.5-Co0.5 with a CsCl type lattice are almost

perfect. [1962Sta] measured at very low temperature (1.3 to 4.2 K) the specific heat of TiCo and Ti4Co3Ni

alloys and found a Debye temperature of 325 and 282 K, respectively.

171


MSIT®

Co–Ni–Ti


Shape memory behavior of the TiNi based alloys is well known. The influence of Co additions has been

investigated by [1998Hos]. In particular, they studied mechanical properties and the Ti and Co

concentration dependence of Ms and As in Ti(Co,Ni) samples with 49 < x(Ti) < 54 and 0 < x(Co) < 6 mol%.

Up to 50 mol% TiCo, the change in hardness is very slight, but with higher cobalt contents the hardness of

the alloys increases and passes through a flat maximum at 80 mol% TiCo [1975Kor].

Miscellaneous

Mutual diffusion coefficients for all the Co-Ni-Ti phases have been calculated by [1982Gry] on the basis of

the diffusion couples results obtained by [1979Sag, 1980Gry].

Samples at the Ti21Co70Ni9 and Ti22Co68Ni10 compositions have been investigated by [1997Liu] while

searching for ternary and quaternary precipitation hardenable alloys: plate shaped gamma precipitates have

been found in aged samples richer in Co.

The partition of Co between (Ni) and TiNi3 phases has been studied by [1994Jia] in the 1000-1200°C

temperature range by the diffusion couple method, as part of a systematic investigation on the partition of

several transition elements between the above mentioned phases.

References

[1962Sta] Starke, E.A., Cheng, C.H., Beck, P.A., “Low Temperature Specific Heat of Ti Alloys with

CsCl-Type Ordered Structure”, Phys. Rev., 126(5), 1746-1748 (1962) (Experimental, 9)

[1966Vuc] Van Vucht, J.H.N., “Influence of Radius Ratio on the Structure of Intermetallic Compounds

of the AB3 Type”, J. Less-Common Met., 11, 308-322 (1966) (Crys. Structure, 21)

[1950Duw] Duwez, P., Taylor, J.L., “The Structure of the Intermediate Phases in Alloys of Titanium

with Iron, Cobalt and Nickel”, Trans. AIME, 188, 1173-1176 (1950) (Crys. Structure,

Experimental)

[1954Poo] Poole, D.M., Hume-Rotary, W., “The Equilibrium Diagram of the System

Nickel-Titanium”, J. Inst. Met., 83, 473-480 (1954-55) (Equi. Diagram, Experimental)

[1959Yur] Yurko, G.A., Barton, J.W., Parr, J.G., “The Crystal Structure of Ti2Ni”, Acta Crystallogr.,

12, 909-911 (1959) (Crys. Structure, Experimental)

[1969Aok] Aoki, Y., Nakamichi, T., Yamamoto, M., “Magnetic Properties of Cobalt-Titanium Alloys

with the CsCl-Type Structure”, J. Phys. Soc. Jpn., 27, 1455-1458 (1969) (Crys. Structure,

Magn Prop., Experimental 13)

[1969Sin] Sinha, A.K., “Close-Packed Ordered AB3 Structures in Ternary Alloys of Certain

Transition Metals”, Trans. Met. Soc. AIME, 245, 911-917 (1969) (Crys. Structure,

Experimental, 16)

[1971Pet] Pet’kov, V.V., Kireev, M.V., “Intermediate Phases in the Titanium-Cobalt System”,

Metallofizika, 33, 107-115, (1971) (Equi. Diagram, Experimental)

[1975Kor] Kornilov, I.I., Kachur, E.V., Belousov, O.K., “Investigation of the TiNi-TiCo System”,

Russ. Metall., 2, 162-163 (1975), translated from Izv. Akad. Nauk SSSR, Met., 2, 209-210

(1975) (Experimental, Equi. Diagram, Mechan. Prop., Magn. Prop., 12)

[1979Sag] Sagyndykov, A.S., Raevskaya, M.V., Sokolovskaya, Y.M., Gryzunov V.U., “Use of a

Method of Diffusion Layers for Plotting the Phase Diagram of the Titanium-Nickel-Cobalt

System”, (in Russian) Kompleksn. Ispol'z. Mineral’n Syr'ya, (4), 47-50 (1979),


[1980Gry] Gryzunov, V.I., Sagyndykov, A.S., “Mutual Diffusion in the System Ti-Ni-Co”, Fiz. Met.

Metalloved., 49(5), 1103-1107 (1980), translated from Phys. Met. Metall., 49(5), 178-182

(1980) (Experimental, Equi. Diagram, Calculation, Phys. Prop., 8)

[1981Loo] Van Loo, F.J.J., Bastin, G.F., “Phase Relations and Diffusion Paths in the Ti-Ni-Co System

at 900°C”, J. Less-Common Met., 81, 61-69 (1981) (Experimental, Equi. Diagram, Crys.

Structure, 11)

172


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Co–Ni–Ti

[1982Gry] Gryzunov, V.I., Sagyndykov, A.S., Sokolovskaya, Ye.M., “Temperature Dependence of

Mutual Diffusion Coefficients in the Ti-Ni-Co System”, Phys. Met. Metallogr., 54(6), 69-73

(1982) (Theory, Phys. Prop., 13)

[1983Gry] Gryzunov, V.I., Shcherbedinskiy, G.V., Sokolovskaya, B.K., “Kinetics of Phase Growth

During Interdiffusion in Ternary Multiphase Metallic Systems”, Fiz. Met. Metalloved.,

56(1), 194-197 (1983), translated from Phys. Met. Metall., 56(1), 183-186 (1983) quoted in

[1993Du]

[1985Shi] Shinoda, T., Isobe, Y., Suzuki, T., “The (fcc)/ (hcp)-Phase Stability in the

Cobalt-Nickel-Titanium System”, Z. Metallkd., 76, 600-608 (1985) (Equi. Diagram,

Thermodyn., Calculation, Experiental, 21)

[1986Kra] Krasheninnikova, N.G., Mogutnov, B.M., Tomilin, I.A., Shaposhnikov, N.G., Erivanskaya

T.Yu., “Specific Heat and Entropy of the Ternary Compounds in the Ni-Co-Ti System at

Low Temperature” (in Russian), Dokl. Akad. Nauk SSSR, 289(5), 1156-1159 (1986)

(Experimental, Calculation, Thermodyn., 5)

[1986Liu] Liu, Y., Takasugi, T., Izumi, O., “Alloying Behavior of Co3Ti”, Met. Trans. A, A17

1433-1439 (1986) (Equi. Diagram, Experimental, Crys. Structure, 21)

[1986Zas] Zasypalov, Yu.V., Kiselev, O.A., Mogutnov, B.M., “Enthaly of Formation of Intermetalic

Compounds CoTi1-xAlx and TiNi1-xCox (0 x 1)”, Dokl. Akad. Nauk SSSR, 287(1), 158

(1986) (Thermodyn., Experimental, 7)

[1987Kra] Krasheninnikova, N.G., Mogutnov, B.M., Tomilin, I.A., Shaposhnikov, N.G.,

“Thermodynamic Properties of the Intermetallic Compounds (Ni0.5Ti0.5)x(Co0.5Ti0.5)1-x

and (Co0.5Ti0.5)x(Co0.5Al0.5)1-x at Low Temperatures”, Russ. Phys. Chem., 61(11),

1627-1630 (1987), translated from Z. Fiz. Khim., 61, 3089-3093 (1987) (Experimental,

Calculation, Thermodyn., 10)

[1988Kra] Krasheninnikova, N.G., Mogutnov, B.M., Tomilin, I.A., Shaposhnikov, N.G.,

Kaloshkin, S.D., “Heat Capacity of Intermetallic Compound Ni0.85Co0.15Ti in the

Temperature Range 78-273 K”, Phys. Met. Metallogr., 65(5), 184-185 (1988)


[1991Mur] Murray, J.L., “The Ni-Ti (Nickel-Titanium) System”, in “Monograph Series on Alloy Phase

Diagrams - Phase Diagram of Binary Nickel Alloys”, Nash, P., (Ed.), ASM International 6,

197-210 (1991) (Equi. Diagram, Crys. Structure, Thermodyn., Review, 110)

[1991Nis] Nishizawa, T., Ishida, K., “Co-Ni (Cobalt-Nickel)” in “Monograph Series on Alloy Phase

Diagrams - Phase Diagram of Binary Nickel Alloys”, Nash, P., (Ed.), ASM International 6,

69-74 (1991) (Equi. Diagram, Crys. Struct., Thermodyn., Review, 91)

[1992Gou] Goubaa, K., Jordan, L., Masse, M., Bouquet, G., “Efficiency of Various Techniques in

Detecting the “R”-Phase in Ni-Ti, Ni-Ti-Cu and Ni-Ti-Co Shape Memory Alloys”, Scr.

Metall. Mater., 26, 1163-1168 (1992) (Experimental, 13)

[1992Shi] Shimizu, K., Tadaki, T., “Recent Studies on the Precise Crystal-Structural Analyses of

Martensitic Transformation”, Mater. Trans., JIM, 33(3), 165-177 (1992) (Calculation, Crys.


[1993Du] Du, Y., Jin, Z.P., Huang, P.Y., “Determination of the 850°C Isothermal Section in the

Co-Ni-Ti System”, J. Phase Equilib., 14(3), 348-353 (1993) (Experimental, Equi.

Diagram, 8)

[1994Jia] Jia, C.C., Ishida, K., Nishizawa, T., “Partitioning of Alloying Elements Between (A1) and

(DO24) Phases in the Ni-Ti Base Systems”, Exp. Methods Phase Diagram Determ., Proc.

Symp., 1993 (Pub. 1994), Morral, J.E., Schiffman, R.S., Merchant, S.M., (Eds.), The

Minerals, Metals & Materials Society, 31-38., (1994) (Experimental, Equi. Diagram, Phys.

Prop., 8)

[1996Bel] Bellen, P., Hari Kumar, K.C., Wollants, P., “Thermodynamic Assessment of the Ni-Ti

Phase Diagram”, Z. Metallkd., 87, 972-978 (1996) (Equi. Diagram, Thermodyn.,

Assessment, 43)

173


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Co–Ni–Ti

[1997Liu] Liu, W., Roesner, H., Nembach, E., “Search for Precipitate Hardenable Quaternary

L12-ordered ’-Intermetallics with Compositions Around Ni3(Al,Si,Ti), (Ni,Co)3(Al,Ti)

and (Ni,Co)3(Si,Ti)”, Z. Metallkd., 88(8), 648-651 (1997) (Experimental, Review, 4)

[1998Hos] Hosoda H., Hanada S., Inoue K., Fukui T., Mishima Y., Suzuki T., “Martensite

Transformation Temperatures and Mechanical Properties of Ternary NiTi Alloys with

Offstoichiometric Compositions”, Intermetallics, 6(4), 291-301 (1998) (Experimental,

Mechan. Prop., 24)

[1999Gup] Gupta, K.P., “The Co-Ni-Ti System (Cobalt-Nickel-Titanium)”, J. Phase Equilib., 20(1),

65-72 (1999) (Review, Thermodyn., 12)

[1999Lek] Lekston, Z., Naish, V.E., Novoselova, T.V., Sagaradze, I.V., “Structure of the Alloy

Ti50Ni48.7Co1.3”, Acta Crystallogr., A55, 803-810 (1999) (Crys. Structure,

Experimental, 13)

[1999Zha] Zhao, J.C., “The fcc/hcp Phase Equilibria and Phase Transformation in Cobalt-Based

Binary Systems”, Z. Metallkd., 90(3), 223-232 (1999) (Assessment, Equi. Diagram, 85)

[2000Cac] Cacciamani, G., Ferro, R., Ansara, I., Dupin, N., “Thermodynamic Modelling of the Co-Ti

System”, Intermetallics, 8, 213-222 (2000), and corrigendum in Intermetallics, 9, 179

(2001) (Thermodyn., Assessment, 42)

[2001Dav] Davydov, A.V., Kattner, U.R., Josell, D., Blendell, J.E., Waterstrat, R.M., Shapiro, A.J.,

Boettinger, W.J., “Determination of the CoTi Congruent Melting Point and

Thermodynamic Reassessment of the Co-Ti System”, Metall. Mater. Trans. A, 32A,

2175-2186 (2001) (Equi. Diagrams, Thermodyn., Assessment, 61)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Co,Ni)

( Co)

1495 - 422

(Ni)

< 1455

cF4

Fm3m

Cu

a = 356.6 0.2

a = 356.88

a = 354.47

a = 352.40

0-100 at.% Ni at 0 at.% Ti and > 422°C

0-~10 at.% Ti at 900°C [1981Loo]

at 54 at.% Co, 7.5 at.% Ti [1981Loo]

at 520°C [V-C2]

[Mas2]

at 25°C [Mas2]

( Co)

< 422

hP2

P63/mmc

Mg

a = 250.71

c = 406.86

at 25°C [Mas2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65

a = 322.4 0.2

a = 321.4

Dissolves

~14 at.% Co at 0 at.% Ni and 1020°C

~10 at.% Ni at 0 at.% Co and 942°C

pure Ti [V-C2]

at 9.5 at.% Ni and 0 at.% Co [1981Loo]

at 12 at.% Co and 0 at.% Ni [1981Loo]

174


MSIT®

Co–Ni–Ti

( Ti)

< 882

hP2

P63/mmc

Mg a = 295.06

c = 468.35

Dissolves < 2 at.% Co and < 1 at.% Ni

at 25°C [Mas2]

( Ti) hP3

P6/mmm

Ti

a = 462.5

c = 281.3

at 25°C, HP 1 atm [Mas2]

Ti(Co1-xNix)3

TiCo3

1181

cP4

Pm3m

AuCu3

a = 360.91

a = 362.0 0.3

a = 362.1

dissolves up to ~23 at.% Ni at 900°C

[1981Loo]. 75.5 to 80.7 at.% Co at 0

at.% Ni [1971Pet]. Dissolves up to ~15

at.% Ni at 1050°C [1986Liu]

at 23Ti-14Ni (at.%) [1986Liu]

at 25Ti-75Co (at.%), [1981Loo]

at 75.5 at.% Co [1971Pet, V-C2]

Melting point of TiCo3 from calculation

of [2001Dav]

Ti(CoxNi1-x)3

TiNi3 < 1380

hP16

P63/mmc

TiNi3 a = 511.5

c = 834.2

a = 509.6

c = 832.2

a = 510.28

c = 827.19

a = 510.3 0.4

c = 831.6 0.8

dissolves up to ~25 at.% Co at 900°C

[1981Loo]

at x = 0.33 [1969Sin]

75 to ~80 at.% Ni at x = 0

[1969Sin]

[V-C2]

at 25 at.% Ti, 75 at.% Ni [1981Loo]

Ti(Co1-xNix)2(cub)

TiCo2

< 1235

cF24

Fd3m

MgCu2

a = 671.6

a = 670.2 0.2

dissolves ~9 at.% Ni at 900°C

[1981Loo]

66.5 to 67 at.% Co

x = 0 [1971Pet]

at 31.5Ti-3Ni (at.%)[1981Loo]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

175


MSIT®

Co–Ni–Ti

Ti(Co1-xNix)2(hex)

TiCo2

1210

hP24

P63/mmc

MgNi2

a = 473.2

c = 1542.7

a = 473.3 0.4

c = 1543 2

dissolves ~3 at.% Ni at 900°C

[1981Loo]

68.75 to 72 at.% Co

x = 0, at 69 at.% Co [1971Pet]

at 29.5Ti-70.5Co (at.%) [1981Loo]

Ti2(Co,Ni)

Ti2Co

< 1058

Ti2Ni

< 985

cF96

Fd3m

Ti2Ni

a = 1131 1

a = 1128.3

a = 1130

a = 1131.9

a = 1127.8 0.1

continuous solid solution at 900°C

at 17 at.% Co [1981Loo]

from 66.7 to 67.1 at.% Co

[1950Duw, V-C2]

[1971Pet]

from 66 up to 67 at.% Ti

[1954Poo]

[1959Yur, V-C2]

Ti(Co,Ni)

TiNi

< 1311

TiCo

< 1503 5

cP2

Pm3m

CsCl a = 300.1 to 299.0

a = 297.9 0.2

a = 300.2

a = 300.5 0.2

a = 299.2 0.3

a = 295.0

a = 297.0

continuous solid solution at 900°C

[1981Loo]

from 50 at.% Ni to 50 at.% Co and at 50

at.% Ti [1975Kor]

at 44.5Ti-37.5Co (at.%) [1981Loo]

51.5 at.% Ti [1971Pet, V-C2]

at 54 at.% Ni [1981Loo]

45 to 51 at.% Ti [2001Dav]

at 44.5 at.% Ti [1969Aok]

at 50 at.% Ti [1969Aok]

TiNi mP4

P21/m

TiNi

a = 289.8 0.1

b = 410.8 0.2

c = 464.6 0.3

= 97.78 0.04

By martensitic transformation. Single

crystal, diffractometer [V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

176


MSIT®

Co–Ni–Ti

R, Ti(CoxNi1-x) hP*

P31m

a = 740.6 0.3

c = 526.8 0.4

By martensitic transformation,

at 1.3Co-50Ti (at.%) [1999Lek]

* , Ti(Co1-xNix)3 hP24

P62m

Co3V

a = 510.8

c = 1250.9

a = 512.0 0.6

c = 1252 0.2

a = 511.4 0.6

c = 1249 2

0.31 x 0.64 [1981Loo]

at x = 0.50 [1969Sin]

at 25.5Ti-(37 to 47)Co (at.%) [1981Loo]

at 23Ti-34Co (at.%) [1981Loo]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

10 20 30 40

1100

1200

1300

1400

1500

1600

Ti 50.00

Co 50.00

Ni 0.00

Ti 50.00

Co 0.00

Ni 50.00Ni, at.%

Te

mp

era

ture

, °C

L

Ti(Co,Ni)

L+Ti(Co,Ni)

1505°C

1311°C

Fig. 1: Co-Ni-Ti.

Pseudobinary section

at 50 at.% Ti

177


MSIT®

Co–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Co

Ni Data / Grid: at.%

Axes: at.%

(Co,Ni)

TiCo3

TiCo2(h)

TiCo2(c)TiCoTi

2Co

TiNi3

TiNi

Ti2Ni

(βTi)

τ

Fig. 2: Co-Ni-Ti.


900°C

178


MSIT®

Co–Si–Ti

Cobalt – Silicon – Titanium

Oksana Bodak, Nathalie Lebrun

Literature Data

Several ternary compounds have been found in the Co-Si-Ti system.

The ternary compound 1 (Ti6Co16Si7) was revealed by [1962Gla]. The lattice parameters were later

confirmed by [1963Spi, 1966Gla]. The 2 (Ti2Co3Si) phase has been detected by [1965Mar, 1966Mar1,

1992Lut] with a MgZn2 type structure. The same type structure has been reported on a Ti33.3Co26.7Si40

alloy [1978Mit] but large discrepancies in the lattice parameters have been observed. Moreover the

composition of the compound given by [1978Mit] (Ti3Co3Si4) is different from those of [1965Mar,

1966Mar1]. A 3 (TiCo3Si2) compound was first suggested by [1966Mar1] with also an hexagonal type

structure. Another composition of 3 has been proposed by [1974Ste] (TiCo4Si3) with a lattice parameter c

four times lower than the one measured by [1966Mar1]. The ternary compound 4 (TiCoSi) was discovered

by [1963Spi]. Its existence was later confirmed by [1969Jei, 1983Szy, 1984Bas, 1992Lut, 1998Lan].

[1978Mit] also detected the existence of a ternary compound with a TiNiSi type structure which

corresponds to the same ternary compound TiCoSi. Another ternary compound called 5 (TiCoSi2) has been

reported in the literature [1966Mar2, 1967Mar, 1992Lut]. [1969Jei] determined the lattice parameters of an

alloy with the composition Ti4Co4Si7 with lattice parameters close to the ones measured for the 5

[1966Mar2, 1967Mar]. The ternary phase 6 with a TiCo2Si composition was found by [1966Mar1,

1973Web]. The ternary phase 7 discovered by [1966Mar1] with a composition of Ti0.75Co0.25Si2, was

recently confirmed by [2001Hu]. [2001Hu] confirmed the existence of all the previous ternary phases and

found two new ternary phases 8 (Ti4CoSi4) and 9 (Ti3Co2Si).

Few experimental data are available on phase equilibria in the Co-Si-Ti ternary system. A eutectic has been

found by [1978Hao] in Ti30Co60Si10, Ti70Co22.5Si7.5, Ti18Co75Si7 and Ti15Co80Si5 alloys. It involves a

silicide compound of type TixSiy and melts at 1135°C. No composition of this eutectic has been measured.

Two isothermal sections are available in the literature: one at 800°C [1966Mar1] and one at 1100°C

[2001Hu]. Moreover, [1992Lut, 1968Mar] investigated the phase equilibria in the TiCo-TiSi system.

Binary Systems

The binary phase diagrams Co-Si and Co-Ti were accepted from [Mas2].

The partial system TiSi-Si has been recently re-investigated by [1998Du] between 1000 and 1500°C using

X-ray diffraction and DTA techniques. Differences have been observed in the melting temperature of the

TiSi2 phase (1488°C instead of 1500°C in [Mas2]) as well as in the eutectic reaction L TiSi+TiSi2 which

is found to be 7°C higher than that reported in [Mas2]. The phase diagram was accepted from [Mas2], except

for the Si rich region for which the partial diagram TiSi-Si phase diagram was accepted from [1998Du].

Solid Phases

Table 1 summarizes the crystal structure data for the unary, binary and ternary phases.

Identical lattice parameters of 1 and 2 have been measured by [1962Gla, 1963Spi, 1966Gla] and

[1965Mar, 1966Mar1], respectively. [1978Mit] found larger lattice parameters. Since [1978Mit] reported a

composition of 2 as Ti3Co3Si4 which is not in agreement with the one suggested by [1965Mar, 1966Mar1,

2001Hu], the lattice parameters were not retained in this assessment. Only the more recent crystallographic

data are reported in Table 1 since no large discrepancies have been observed in the values of the lattice

parameters of these two ternary compounds. Two possible compositions of the 3 phase have been reported

in the literature: TiCo3Si2 [1966Mar1] and TiCo4Si3 [1974Ste]. No definitive conclusion can be made on

the composition and the value of the lattice parameter c. Consequently, the different literature data have

been included in Table 1. The lattice parameters of the ternary compound 4 have been well defined by

[1969Jei, 1983Szy, 1984Bas, 1998Lan]. [1963Spi] reported larger structural parameters for 4 which were

179


MSIT®

Co–Si–Ti

not retained in this assessment since low resolution technique has been employed (Debye-Scherrer X-ray

photographs). Only the results of [1998Lan] are reported in Table 1 because of the very precise study of the

crystal structure using a Guinier diffractometer and Rietveld analysis of the diffractograms. Two

compositions have been proposed for the 5 phase: TiCoSi2 [1966Mar2, 1967Mar] and Ti4Co4Si7 [1969Jei]

with quasi identical parameters. The ternary phase called 6 (TiCo2Si) has been found by [1966Mar1,

1973Web]. The more recent crystallographic results are reported in Table 1. The crystal parameters of the

ternary compounds 7 and 9 have been recently determined [2001Hu] with an orthorhombic and a

hexagonal structure, respectively. The last ternary compound 8 was recently discovered [2001Hu] but its

structure is unknown until now.

All these ternary phases have been considered as stoichiometric at 800°C, whereas homogeneity ranges

measured at 1100°C show large extensions, especially for 2 and 5.


On the basis of the phase compositions of alloys, authors [1966Mar1] assume presence of pseudobinary

sections CoSi2-TiSi2, Co2Si-Ti5Si3, Co-TiCoSi, TiCo2-TiCoSi, TiCo-Co2Si, TiCo-CoSi.

Isothermal Sections

Two isothermal sections are available in the literature: at 800°C [1966Mar1, 1968Mar] and 1100°C

[2001Hu]. The modified isothermal sections have been reported in Figs. 1 and 2.

At 800°C, most of the ternary phases have been considered as stoichiometric since no experimental

evidence is presented by [1966Mar1, 1968Mar]. The measured solubility range of the ternary phase 2 has

been indicated in the isothermal section since [1966Mar1] mentioned that 2 contains between 12.5 and 20

at.% Si. The solubility ranges of the binary phases at the binary edge were modified in agreement with the

accepted binary phase diagrams. No large solubility range of the binary phases into the ternary have been

observed at 800°C. The maximum solubility range is measured for the Co2Si and the Ti2Co phases with a

maximum solubility of 6 at.% Ti and 5 at.% Si, respectively. At 800°C, [1966Mar1, 1968Mar] did not

mentioned the existence of the binary phases TiCo3, ( Co), Ti2Si and Ti5Si4. Moreover the solubility ranges

of the phases ( Ti) and ( Ti) are not in agreement with the binary phase diagrams accepted here.

Consequently the missing binary phases have been indicated on the ternary section and the solubility ranges

of the Ti unary phases have been modified. Because of the lack of information on the phase equilibria

involving these phases, no phase fields have been reported in the isothermal section.

The solubility ranges of the ternary phases were well measured at 1100°C [2001Hu]. Since no evidence of

solubility range has been reported in [2001Hu], the 8 ternary compound has been considered as

stoichiometric. Apart from 6 phase, the other ternary compounds show considerable homogeneity ranges,

especially the 5 and 2 phases. As reported in the isothermal section at 800°C, the 2 region shows the

shape of a bar that is parallel to the Co-Si side. Two new phases with the compositions Ti4CoSi4 and

Ti3Co2Si have been discovered at 1100°C [2001Hu]. Two binary phases Ti2Si and TiCo3 were not found

at 1100°C by [2001Hu] and [1968Mar] despite they do exist at 1100°C in the binary systems. These solid

phases have been indicated in the isothermal section (Fig. 2). The solubility of Co in Ti5Si3 is quite large

(6 at.% Co). A considerable amount of Si can also be dissolved in the TiCo2 (h) phase. In contrast the Si

solubility in TiCo3 is quite small. All the binary intermetallic phases on the Co-Si side show large

composition ranges. The extensions of the homogeneity ranges of many phases as determined by [2001Hu]

require more a precise determination.

[1978Hao] mentioned the existence at 1135°C of a eutectic transformation in cobalt rich alloys with

unknown structure. The phase structure of alloys on the TiCo-TiSi section at 800°C was determined in

[1992Lut] and they are in general agreement with [1966Mar1, 1968Mar].


The TiCo2Si compound is found to be ferromagnetic with a Curie temperature of 102 4°C [1973Web].

Magnetic and magneto-optical properties of the TiCo2Si phase was measured in [1983Bus]. Magnetic

180


MSIT®

Co–Si–Ti

properties of the TiCo2Si compound were studied in [1973Web, 1983Bus]. Chemical bonds in compounds

with the TiNiSi structure type were analyzed in [1998Lan].

References

[1962Gla] Gladyshevsky, E.I., Markiv, V.Ya., Kuzma, Yu.B., “New Ternary Compounds with a

Structure of the Mg6Cu16Si7 Type” (in Ukrainian), Dop. Akad. Nauk Ukr. SSR, (4), 481-483


[1963Spi] Spiegel, F. X., Bardos, D., Beck, P.A., “Ternary G and E Silicides and Germanides of

Transition Elements”, Trans. Met. Soc. AIME, 227, 575-579 (1963) (Crys. Structure,

Experimental, 13)

[1965Mar] Markiv, V.Ya., Voroshilov, Yu.V., Gladyshevsky, E.I., “Ternary Laves Phases in the

Systems Ti-Co-Si (Ge) and Zr -Fe-Si (Ge)”, Inorg. Mat., 1, 818-821 (1966), translated from

Izv. Akad. Nauk SSSR, Neorg. Mater., 1, 890 (1965) (Crys. Structure, Experimental, 5)

[1966Gla] Gladyshevsky, E.I., Markiv, V.Y., Kus’ma, Y.B., Cherkashin, E.E., Titan i Ego Splavy,

Moscow (in Russian), 10, 73 (1966) cited in [2001Hu]

[1966Mar1] Markiv, V.Ya., Gladyshevsky, E.I., Fedoruk, T.I., “Phase Equilibria in the Ti-Co-Si

System.” (in Russian), Izv. Akad. Nauk SSSR, Met., (3), 179-182 (1966) (Crys. Structure,

Equi. Diagram, Experimental, *, #, 13)

[1966Mar2] Markiv, V.Ya., “The Crystal Structure of the Compounds R(M,X)2 and RMX2 in Zr-Ni-Al,

Ti-Fe-Si and Related Systems”, Acta Crystallogr., 21(7), A84 (1966) (Abstract)

[1967Mar] Markiv, V.Ya., Gladyshevsky, E.I., Skolozdra, R.V., Kripyakevich, P.I. “Ternary

Compounds of the RX’X’’2 Type in the Ti-V(Fe, Co, N i)-Si and Similar Systems”, (in

Russian) Dop. Akad. Nauk Ukr. RSR, A3, 266-268 (1967) (Crys. Structure,

Experimental, 12)

[1968Mar] Markiv, V.Ya., Gladyshevsky, E.I., Kvipyakevich, P.I., Fedoruk, T.I., Lysenko, L.A., “A

Study of the Phase Equilibria and Crystal Structures of Compounds in the Ti-Co-Si System”

(in Russian), Diagrammy Sostoyaniya Metallich. Sistem, Nauka, Moscow, 137-145 (1968)

(Equi. Diagram, Crys. Structure, Experimental, Review, #, *, 29)

[1969Jei] Jeitschko, W., Jordan, A.G., Beck, P.A., “V and E Phases in Ternary Systems with

Transition Metals and Silicon or Germanium”, Trans. Met. Soc. AIME, 245, 335-339 (1969)


[1973Web] Webster, P.J., Ziebeck, K.R.A., “Magnetic and Chemical Order in Heusler Alloys

Containing Cobalt and Titanium”, J. Phys. Chem. Solids, 34, 1647-1654 (1973) (Crys.

Structure, Magn. Prop., Experimental, 26)

[1974Ste] Steinmetz, J., Albrecht, J.M., Malaman, B., “A New Family of Ternary Silicides of the

General Formula TT’4Si3 (T = Ni, Nd, Ta; T’ = Fe, Co, Ni)” (in French), Compt. Rend.

Acad. Sci. Paris, 279C, 1119-1120- (1974) (Crys. Structure, Experimental, 4)

[1978Mit] Mittal, R.C., Si, S.K., Gupta K.P., “Si-Stabilised C14 Laves Phases in the Transition Metal

Systems”, J. Less-Comm. Met., 60, 75-82 (1978) (Crys. Structure, Experimental, 12)

[1978Hao] Haour, G., Mollard, F., Lux, B., Wright, I.G., “New Eutectics Based on Fe, Co, Ni. II.

Co-Base Eutectics”, Z. Metallkd., 69-74 (1978) (Experimental, 7)

[1983Bus] Buschow, K.H.J., Engen, van P.G., Jongebreur, R., “Magneto-Optical Properties of

Metallic Ferromagnetic Materials”, J. Magn. Magn. Mater., 38, 1-22 (1983) (Experimental,

Magn. Prop., 23)

[1983Szy] Szytula, A., Bazela, W., Radenkovic, S., “Crystal and Magnetic Structure of CoMn1-xTixSi

System” J. Magn. Mag. Mater., 38, 99-104 (1983) (Crys. Structure, Experimental, Magn.

Prop., 12)

[1984Bas] Bazela-Wrobel, W., Szytula, A., Leciejewicz, J., “Magnetic Properties of RhMnSi and

CoSiTi”, Phys. Status Solidi A, 82A, 195 (1984) (Crys. Structure, Experimental, Magn.

Prop., 16)

181


MSIT®

Co–Si–Ti

[1992Lut] Lutskaya, N.V., Alisova, S.P., “The Phase Structure of TiCu(TiNi, TiCo)-TiSi Sections in

Ternary TiCu(Ni, Co)-Si Systems”, Russ. Metall. (Engl. Transl.), (3), 180-182 (1992),

translated from Izv. Akad. Nauk SSSR Met., (3), 1992, 194-196 (Equi. Diagram,

Experimental, 9)

[1998Du] Du, Y., Schuster, J.C., “A Re-Investigation of the Constitution of the Partial System

TiSi-Si”, J. Mat. Sci. Lett., 17, 1407-1408 (1998) (Equi. Diagram, Experimental, 7)

[1998Lan] Landrum, G.A., Hoffmann, R., Evers, J., Boysen, H., “The TiNiSi Family of Compounds:

Structure and Bonding”, Inorg. Chem., 37(22), 5754-5763 (1998) (Crys. Structure,

Experimental, 34)

[2001Hu] Hu, X., Chen, G., Cinca, I., “The 1100°C Isothermal Section of the Ti-Co-Si Ternary

System”, J. Phase Equilib., 22(2), 114-121 (2001) (Crys. Structure, Experimental, *, #, 49)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Co)

< 1495 - 422

cF4

Fm3m

Cu

a = 356.3 0 to 16.4 at.% Si

0 to 14.4 at.% Ti

[Mas2, V-C2]

( Co)

< 1250

hP2

P63/mmc

Mg

a = 250.6 0 to 18.4 at.% Si

0 to1 at.% Ti

[Mas2, V-C2]

(Si)

< 1414

cF8

Fd3m

C (diamond)

a = 543.09 [Mas2, V-C2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 331.12 0 to 14.5 at.% Co

0 to 3.5 at.% Si

[Mas2, V-C2]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.03 0 to 0.8 at.% Co

0 to 0.5 at.% Si

[Mas2, V-C2]

Co3Si

1214 - 1193

hP8

P63/mmc

Ni3Sn

a = 497.6 0.2

c = 406.9 0.6

[Mas2, V-C2]

Co2Si

< 1320

oP12

Pnma

Co2Si

a = 491.9

b = 372.5

c = 710.4

32 to 34 at.% Si

[Mas2, V-C2]

dissolves ~6 at.% Ti

[1966Mar1, 1968Mar]

Co2Si

1334 - 1238

- - 32 to 35.8 at.% Si

[Mas2, V-C2]

CoSi

< 1460

cP8

P213

FeSi

a = 445.0 49 to 52 at.% Si

[Mas2, V-C2]

CoSi2< 1326

cF12

Fm3m

CaF2

a = 535 [Mas2, V-C2]

dissolves ~2 at.% Ti

[1966Mar1]

182


MSIT®

Co–Si–Ti

TiCo3

< 1190

cP4

Pm3m

AuCu3

a = 362.8 19.3 to 24.5 at.% Ti

[Mas2, V-C2]

TiCo2 (h)

< 1210

hP24

P63/mmc

MgNi2

a = 473

c =1541

28 to 31.25 at.% Ti

[Mas2, V-C2]

TiCo2 (c)

< 1235

cF24

Fd3m

MgCu2

a = 669.2 33 to 33.5 at.% Ti

[Mas2, V-C2]

dissolves ~2 at.% Si

[1966Mar1]

TiCo

< 1325

cP2

Pm3m

CsCl

a = 300.2 45 to 51 at.% Ti

[Mas2, V-C2]


[1966Mar1]

Ti2Co (c)

< 1058

cF96

Fd3m

NiTi2

a = 1129.5 66.7 to 67.1 at.% Ti

[Mas2, V-C2]


[1966Mar1]

Ti3Si

< 1170

tP32

P42/m

Ti3P

a = 1020.6 0.6

c = 506.9 0.2

[Mas2, V-C2]

Ti5Si3< 2130

hP16

P63/mcm

Mn5Si3

a = 746.10 0.03

c = 515.08 0.01

35.5 to 39.5 at.% Si

[Mas2, V-C2]

dissolves ~3 at.% Co

[1966Mar1]

Ti5Si4< 1920

tP36

P41212

Zr5Si4

a = 713.3

c = 1297.7

[Mas2, V-C2]

TiSi

< 1570

oP8

Pnma

FeB

a = 655.1 0.6

b = 363.3 0.3

c = 498.3 0.5

[V-C2]


[1966Mar1]

TiSi2< 1488

oF24

Fddd

TiSi2

a = 826.71 0.09

b = 480.00 0.05

c = 855.05 0.11

[Mas2, V-C2]


[1966Mar1]

* 1, Ti6Co16Si7 cF116

Fm3m

Mg6Cu16Si7 or

Mn23Th6

a = 1123.2 0.4 [1966Gla]

* 2, Ti2Co3Si hP12

P63/mmc

MgZn2

a = 479.7 0.3

c = 756.4 0.3

[1965Mar, 1966Mar1]

* 3, TiCo3Si2

TiCo4Si3

hP*

P6/mmm a = 1697

c = 3179

a = 1700.1

c = 795.0

Ti14Co49Si37 in [2001Hu]

[1966Mar1, V-C2]

[1974Ste, V-C2]; dm = 6.44

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

183


MSIT®

Co–Si–Ti

* 4, TiCoSi oP12

Pnma

TiNiSi or CoSi2

a = 610.0 0.2

b = 371.5 0.1

c = 692.7 0.2

[1998Lan, V-C2]

* 5,

TiCoSi2

Ti4Co4Si7

t**

TiNiSi2tI56

I4/mmm

Zr4Co4Ge7

a = 1247

c = 493

a = 1251.3 0.3

c = 493.4 0.1

[1966Mar2]

[1969Jei, V-C2]

* 6, TiCo2Si cF16

Fm3m

BiF3

a = 574.0

L21 structure type

[1973Web]

* 7, Ti3CoSi8 o** a = 796.1

b = 704.8

c = 546.7

[2001Hu]

* 8, Ti4CoSi4 Structure unknown [2001Hu]

* 9, Ti3Co2Si t** a = 673.5

c = 978.1

[2001Hu]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

20

40

60

80

20 40 60 80

20

40

60

80

Ti Co


Axes: at.%

τ7

τ5

τ4

τ3

τ1

τ6

τ2

(αCo)

(εCo)

αCo2Si

CoSi

CoSi2

(Si)

TiSi2

TiSi

Ti5Si

4

Ti5Si

3

Ti2Si

(αTi)

(βTi) Ti2Co TiCo TiCo

2(c)

TiCo2(h)

TiCo3

Fig. 1: Co-Si-Ti.


800°C

184


MSIT®

Co–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Co


Axes: at.%(Si)

CoSi2

CoSi

αCoSi2

(εCo)

(αCo)

TiCo3TiCo

2(h)TiCo

2(c)TiCo

L(βTi)

Ti5Si

3

Ti5Si

4

TiSi

TiSi2 τ

7

τ5

τ3

τ4

τ1τ6τ

2τ

9

τ8

Ti2Si

Fig. 2: Co-Si-Ti.


1100°C

185


MSIT®

Cr–Nb–Ti

Chromium – Niobium – Titanium

Gautam Ghosh

Literature Data

A fairly large number of experimental studies have been carried out to establish the ternary phase equilibria

[1962Sha, 1962Sve1, 1962Sve2, 1963Sha1, 1963Kor, 1964Koc, 1964Sve, 1965Kor, 2002Tho]. Some of

these results were reviewed by [1973Bud]. [1962Sha] determined three partial isothermal sections at 600,

800 and 1000°C. [1962Sve1] determined and isothermal section at 1250°C, and an isopleth along

NbCr2-TiCr2 and another isopleth at 10 mass% Ti. The first comprehensive study of phase equilibria was

carried out by [1965Kor]. They prepared a large number of alloys using iodide grade Ti, 99.27% Nb and

99.98% Cr. The alloys were prepared by both arc-melting in an argon atmosphere and levitation melting in

a helium atmosphere. The alloys were homogenized in the temperature range 1300 to 1500°C for up to

240 h depending on the alloy composition. For determining isothermal sections in the temperature range of

600 to 1000°C, the alloys were annealed further for up to 550 h. Conventional metallography and X-ray

diffraction were used to establish the phase equilibria. The results were presented in terms seven isothermal

sections, from 1900 to 1300°C, and four vertical sections.

[2001Yos] determined the phase equilibria of the Cr corner at 1250°C using optical microscopy, X-ray

diffraction and analytical electron microscopy. They prepared ten ternary alloys using 99.9% Cr, 99.5% Nb

and 99.9% Ti by arc melting, and subsequently they were annealed at 1250°C for 24 h. The tie lines between

the phases were established by quantitative analytical electron microscopy. Besides experimental phase

equilibria studies, there are two reports on thermodynamic calculation of phase equilibria using CALPHAD

(Calculation of Phase Diagrams) methodology [1975Kau, 2000Lee, 2001Kau].

[2004Zha] investigated solid-solid phase equilibria at 1000, 1150 and 1200°C using diffusion multiples.

They also reported three vertical sections along Nb:Ti=1:3, Nb:Ti=1:1 and Nb:Ti=3:1 showing the

solubility of Cr in (Nb,Ti), with respect to C15 Laves phase, in the temperature range of 800 to 1600°C.

The diffusion multiples were prepared using high purity Cr, Nb and Ti which were subjected to hot isostatic

pressing at 1204°C and 200 MPa for 4 h. The entire assembly was then encapsulated in quartz tubes,

containing yttrium in tantalum foil as getters for interstitials (C, N and O), backfilled with pure Ar. The

encapsulated samples were then annealed at 1000°C (for 4000 h), 1150°C (for 2000 h) and 1200°C (for

1000 h). The composition of phases in the interdiffusion zone was measured by quantitative electron probe

microanalysis technique, and the structural information of the phases was obtained by electron

backscattered diffraction analysis.

Binary Systems

The Cr-Nb, Cr-Ti and Nb-Ti binary phase diagrams are accepted from [2004Iva1], [2004Iva2] and

[2001Zha], respectively. In the Cr-Nb equilibrium diagram there are two Laves phases NbCr2 (C36) and

NbCr2 (C15), while all three Laves phases TiCr2 (C14), TiCr2 (C36) and TiCr2 (C15) are stable in the

Cr-Ti equilibrium phase diagram. In the case of NbCr2, the polymorphic transformation C15 C36 was

assumed to be a first-order. In the ternary system, the polymorphic transformations of the Laves phases are

also assumed to be first-order.

Solid Phases

There is no ternary phase in this system. The details of crystal structures and lattice parameters of the solid

phases are listed in Table 1.

The stability of Laves phases has been discussed a number of times [1997Zhu, 1998Tak, 2002Tho].

[1997Zhu] suggested that the average valence electron concentration (e/a) is a dominant factor in

controlling the stability of NbCr2-based transition-metal Laves phases. They proposed the following

186


MSIT®

Cr–Nb–Ti

empirical rule: C14 is stabilized in the e/a range of 5.88 to 7.53; C15 is stabilized when e/a 5.76 and e/a

7.65; C36 is stabilized when 5.88 > e/a > 5.76 and 7.65 > e/a > 7.53.

The site occupancy of Ti in C15- NbCr2 has been studied in detail [1998Kot, 1999Oka] using ALCHEMI

(Atom Location by CHanneling Enhanced MIcroanalysis) technique in a transmission electron microscope.

[1998Kot] used Cr68Nb15Ti17 alloy while [1999Oka] used Cr66.7-x/2Nb33.3-x/2Tix and Cr66.7Nb33.3-xTixalloys. In all cases Ti prefers to occupy Nb sublattice, while Nb partitions to both sublattices. These results

has been discussed both in terms of size effect [1998Kot] and electronic structure viewpoint [2002Tho].

At 1200°C, NbCr2 and TiCr2 phases form a continuous solid solution. The lattice parameter of

(Nb33-xTix)Cr67 shows a significant negative departure from a linear rule of mixtures suggesting tighter

binding than binary alloys [2002Tho]. The authors [1998Che, 2002Tho] reported the lattice parameter of

C15-(Nb33-xTix)Cr67 as a function of composition. The upper and lower limits are given in Table 1. The

variation of lattice parameter of the solid solution between NbCr2 and TiCr2 was also reported by

[1962Sha] in alloys quenched from 600°C corresponding to the two-phase field C15+bcc. They agree fairly

well with more recent lattice parameter data of Nb1-xTixCr2 [1998Che, 2002Tho]; however, [1962Sha] did

not determine the composition of Nb1-xTixCr2 phase.

[1995Tho] obtained a metastable bcc phase in alloys along NbCr2-TiCr2 and Ti-NbCr2 sections where the

formation of Laves phases were suppressed by splat quenching. The lattice parameter of bcc phase is shown

to obey Vegard’s law.


[1962Sve1] reported approximate liquidus and solidus isotherms for the composition range

Cr-NbCr2-TiCr2. [1963Sha2] determined the solidification temperature of several ternary alloys. They

reported the solidification temperature in both tabulated and graphical forms; however, it is not clear if the

solidification temperature referred to liquidus or solidus. Their graphical plot “solidification temperature”

showed significant discrepancy with the accepted binary phase diagrams. [1964Koc] also reported liquidus

and solidus isotherms for the entire composition range. Once again, these isotherms also show significant

discrepancy with the accepted binary phase diagrams.

Isothermal Sections

Figures 1 to 7 show the isothermal sections from 1900 to 1300°C, at 100°C interval [1965Kor]. Most of

these are constructed from the results of vertical sections. The isothermal section at 1600°C proposed by

[1965Kor] is inconsistent with thermodynamic principles. Consequently, the liquidus shape has been

changed (see Fig. 4). Recently, [1992Tho] calculated the phase diagram at 1400°C which is in good

agreement with the experimental isothermal section reported by [1965Kor]. The isothermal section shown

in Fig. 6 is a compilation of phase diagrams reported by [1965Kor] and [1992Tho].

In Fig. 7, TiCr2 Laves phase should be stable at the Cr-Ti binary edge at 1300°C. However, the

composition trajectory for TiCr2 -> TiCr2 transformation is not known, and hence it is shown dotted.

Furthermore, additional phase fields, such as ( Ti)+ TiCr2 and ( Ti)+ TiCr2+ TiCr2 are expected to be

present very close to the Cr-Ti edge. Figure 8 shows the isothermal section at 1250°C adopted from

[1962Sve1]; however, the phase equilibria involving (Cr), TiCr2 and NbCr2 are taken from recent results

of [2001Yos]. Results of [1965Kor] and [2001Yos] show that Nb stabilizes NbCr2 phase.

Figures 9, 10, 11, 12, 13 and 14 show the isothermal sections at 1200°C [2004Zha], 1150°C [2004Zha],

1000°C [1962Sha, 2004Zha], 950°C [2002Tho], 800°C [1962Sha] and 600°C [1962Sha], respectively.

Both Nb1-xTixCr2 and TiCr2 Laves phases are stable at the Cr-Ti binary edge in the temperature range of

950 to 1200°C as seen in Figs. 9 to 12. This inevitably causes the presence of a three-phase field

(Cr)+ Nb1-xTixCr2+ TiCr2 shown with dashed lines, as it has not been experimentally verified. However,

it is important to note that this situation is similar to the case of Cr-Ti-V system [2002Gho]. In Figs. 1 to 14,

several adjustments were made to comply with the accepted binary phase diagrams.

187


MSIT®

Cr–Nb–Ti


Several temperature-composition sections were determined [1962Sve1, 1964Koc, 1964Sve, 1965Kor]. It is

interesting to note that NbCr2 and TiCr2 do not form a pseudobinary section. Due to the disagreement

between the vertical section at a constant Cr-content of 66.7 at.% and the accepted Cr-Ti binary phase

diagram, the isopleth proposed by [1962Sve1] and [1964Sve] is not considered in this evaluation. Figure 15

shows the polythermal section along Ti-NbCr2 [1964Sve]. Isopleths at constant mass ratios of 1:4, 2:3, 3:2

and 4:1 were reported by [1965Kor]. [1964Koc] determined eight isopleths at a constant Ti-content of 5,

10, 15, 20, 25, 30, 35, 40 and 70 mass%. [2004Zha] reported three vertical sections along Nb:Ti=1:3,

Nb:Ti=1:1 and Nb:Ti=3:1 (atomic ratios) showing the solubility of Cr in (Nb,Ti) in the temperature range

of 800 to 1600°C. Their measured solubility agrees fairly well with those reported by [1962Sha]. Selected

isopleths are shown in Figs. 15, 16, 17, 18, 19, 20, and 21. Like isothermal sections, adjustments were made

in the temperature-composition sections to comply with the accepted binary phase diagrams.

Thermodynamics

There is no measured thermodynamic data for the ternary alloys. [1975Kau] employed the CALPHAD

technique to calculate isothermal sections at 1300, 1500, 1700 and 1900°C which were in good agreement

with the experimental results of [1965Kor]. [1975Kau] used only the binary interaction parameters. Later,

[2001Kau] calculated isothermal sections at 1300, 1500, 1600 and 1800°C by considering ternary solubility

of the Laves phases. [2000Lee] also employed the CALPHAD technique to derive an optimized set of

ternary interaction parameters for the bcc phase using experimental phase diagram of [1962Sha]. They also

reported calculated isothermal sections at 800 and 1000°C which were in good accord with the experimental

data.


The ambient temperature elastic properties (bulk, shear and Young’s moduli, and Poisson’s ratio), hardness

and indentation fracture toughness of the C15 Laves phase ( NbCr2- TiCr2) were measured by [2002Tho].

[1998Che] also reported hardness and fracture toughness of single phase (C15) (Nb,Ti)Cr2 alloys. The

elastic moduli and hardness generally decreases along the constant Nb/Ti ratio [2002Tho]. With the

substitution of Nb by Ti in NbCr2, the shear moduli and hardness showed a positive deviation with respect

to a linear rule of mixture between NbCr2 and TiCr2. However, the toughness increased only along the

constant Nb/Ti ratio. [1998Che] found that the substitution of Ti by Nb causes an increase in hardness and

a decrease in fracture toughness of the C15 phase. The Vickers hardness values range from 871 to 914

kg mm-2 [1998Che], and 840 to 890 kg mm-2 [2002Tho], while the indentation fracture toughness values

range from 0.69 to 0.82 MPa m0.5 [1998Che] and 1.1 to 1.24 MPa m0.5 [2002Tho].

[1965Sha] determined the elastic (Young’s and shear moduli) and plastic (hardness) properties of single-

and two-phase alloys as a function heat treatment. The composition of single-phase alloys range from

Ti-rich to Nb-rich. They found that Cr is a better solid solution strengthener than Ti. [1962Sve2] measured

hardness of Cr-rich and Ti-rich alloys as a function of temperature (from 20 to 1000°C); however, the alloys

were single phase (bcc, hcp, C15), two-phase (bcc+C15) and three-phase (bcc+hcp+C15).

Fracture toughness and fatigue crack growth resistance were measured by [1996Dav] where Ti was found

to increase the toughness of solid solution of Cr-Nb alloys. [1997Cha] studied the fracture and fatigue

behavior of in situ composites based on the Cr-Nb-Ti ternary system and showed an increase in fracture

resistance with a decreasing volume fraction of NbCr2 particles.

188


MSIT®

Cr–Nb–Ti

References

[1962Sha] Shakhova, K.I., Budberg, P.B., “Phase Diagram of the Ti-Nb-Cr System” (in Russian),

Russ. Metall. Fuels, (6), 72-78 (1962), translated from Izv. Akad. Nauk SSSR, (6), 137-141

(1962) (Equi. Diagram, Experimental, #, *, 5)

[1962Sve1] Svechnikov, V.N., Kocherzhinsky, Yu.A., Latysheva, V.I., Pan, V.M., “The Cr-Nb-Ti

System” (in Russian), Sb. Nauchn. Tr. Inst. Metallofiz., (16), 128 (1962) (Equi. Diagram,


[1962Sve2] Svechnikov, V.N., Kocherzhinsky, Yu.A., Latysheva, V.I., Pan, V.M., “Investigation of

Alloys in the Cr-Nb-Ti System” (in Russian), Issled. po Zharoprochn. Splav., A, (8), 56-61

(1962) (Equi. Diagram, Experimental, #, *, 9)

[1963Sha1] Shakhova, K.I., Budberg, P.B., “Ternary Alloys of the Ti-Nb-Cr System” (in Russian),

Titan i ego Splavy, (10), 37-41 (1963) (Equi. Diagram, Experimental, #, *, 5)

[1963Sha2] Shakhova, K.I., Budberg, P.B., “Solidification Titanium-Chromium-Niobium Alloys”,

Russ. Metall. Mining, (10), 118-119 (1963) (Equi. Diagram, Experimental)

[1963Kor] Kornilov, I.I., Shakhova, K.I., Budberg, P.B., Nedumov N.A., “The Equilibrium Diagram

of TiCr2-NbCr2” (in Russian), Dokl. Akad. Nauk SSSR, 149(6), 1340-1342 (1963) (Equi.

Diagram, Experimental, #, *, 7)

[1964Koc] Kocherzhinsky, Yu.A., Latysheva, V.I., “Solubility in the System Cr-Nb-Ti” (in Russian),

Sb. Nauchn. Tr. Inst. Metallofiz., (20), 125-129 (1964) (Equi. Diagram, Experimental,

#, *, 13)

[1964Sve] Svechnikov, V.N., Kocherzhinsky Yu.A., Latysheva V.I., “Phase Diagrams of the

NbCr2-TiCr2 and NbCr2-Ti Systems” (in Russian), Sb. Nauchn. Tr. Inst. Metallofiz., (19),

192-195 (1964) (Equi. Diagram, Experimental, #, *, 3)

[1965Kor] Kornilov, I.I., Shakhova, K.I., Budberg, P.B., “Phase Equilibrium Diagram of the Ti-Nb-Cr

System”, Russ. Metall., (4), 119-127 (1965) (Equi. Diagram, Experimental, #, *, 5)

[1965Sha] Shakhova, K.I., Budberg, P.B., “Certain Mechanical Properties of Alloys in the Ti-Nb-Cr

System”, Russ. Metall., (2), 66-73 (1965), transl. from Izv. Akad. Nauk SSSR, Met., (2),

1965, 128 (Equi. Diagram, Experimental, #, *, 14)

[1973Bud] Budberg, P.B., “Phase Diagrams of Ternary Systems of Titanium and Chromium with

Niobium, Tantalum and Vanadium”, Khim. Metal. Splavov, Publ. Nauka, Moskow, 85-89

(1973) (Equi. Diagram, Review, 12)

[1975Kau] Kaufman, L., Nesor, H., “Calculation of Superalloy Phase Diagrams: Part III”, Metall.

Trans., 6, 2115-2122 (1975) (Equi. Diagram, Thermodyn., *, 35)

[1992Tho] Thoma, D.J., PhD. Thesis, Univ. Wisconsin, Madison, WI (1992) as quoted in [1996Dav]

[1995Tho] Thoma, D.J., Perepezko, J.H., “Metastable BCC Phase Formation in Nb-Cr-Ti System”,

Mater. Sci. Forum, 19-181, 769-774 (1995) (Crystal Structure, Experimental, Equi.

Diagram, 16)

[1996Dav] Davidson, D.L., Chan, K.S., Anton, D.L., “The Effects on Fracture Touhhness of Ductile

Phase Composition and Morphology in Nb-Cr-Ti and Nb-Si in Situ Composites”, Metall.

Mat. Trans. A, 27, 3007-3018 (1996) (Experimental, Mechan. Prop., 20)

[1997Cha] Chan, K.S., Davidson, D.L., Anton, D.L., “Fracture Toughness and Fatigue Crack Growth

in Rapidly Quenched Nb-Cr-Ti in Situ Composites”, Metall. Mat. Trans. A, 28, 1797-1808


[1997Zhu] Zhu, J.H., Liaw, P.K., Liu, C.T., “Effect of Electron Concentration on the Phase Stability of

NbCr2-Based Laves Phase Alloys”, Mater. Sci. Eng. A, 239-240, 260-264 (1997) (Crys.

Structure, Review, 30)

[1998Kot] Kotula, P.G., Carter, C.B., Chen, K.C., Thoma, D.J., Chu, F., Mitchell, T.E., “Defects and

Site Occupancies in Nb-Cr-Ti C15 Laves Phase Alloys”, Scr. Mater., 39 (4/5), 619-623


189


MSIT®

Cr–Nb–Ti

[1998Che] Chen, K.C., Allen, S.M., Livingston, J.D., “Factors Affecting the Room-Temperature

Mechanical Properties of TiCr2-Base Phase Alloys”, Mater. Sci. Eng. A, A242, 162-173

(1998) (Crys. Structure, Experimental, Mechan. Prop., 51)

[1998Tak] Takasugi, T., Yoshida, M., “The Effect of Ternary Addition on Structure and Stability of

NbCr2 Laves Phases”, J. Mater. Res., 13(9), 2505-2513 (1998) (Crys. Structure,

Experimental, 28)

[1999Oka] Okaniwa, H., Shindo, D., Yoshida, M., Takasugi, T., “Determination of Site Occupancy of

Additives X (X=V,Mo,W and Ti) in the Nb-Cr-X Laves Phase by Alchemi”, Acta Mater.,

47 (6), 1987-1992 (1999) (Crys. Structure, Experimental, 11)

[2000Lee] Lee, J.Y., Kim, J.H., Lee, H.M., “Effect of Mo and Nb on the Phase Equilibrium of the

Ti-Cr-V Ternary System in the Non-Burning -Ti Alloy Region”, J. Alloys Compd., 297,

231-239 (2000) (Equi. Diagram, Thermodyn., *, 19)

[2001Kau] Kaufman, L., “Calculation of Multicomponent Phase Diagrams for Niobium Alloys” in

“Niobium Science and Technology”, Proc. Int. Symp. Niobium, TMS, Orlando, Florida,

107-120 (2001) (Calculation, Equi. Diagram, 20)

[2001Yos] Yoshiba, M., Yaegashi, T., Murakami, Y., Shindo, D., Takasugi, T., “Evaluation of

Microstructures of Nb-Cr-Ti Alloy System by Means of Analytical Transmission Electron

Microscopy” (in Japanese), J. Jpn. Inst. Met., 65(5), 389-396 (2001) (Crys. Structure,

Experimental, Equi. Diagram, #, *, 25)

[2001Zha] Zhang, Y., Liu, H., Zhanpeng, J., “Thermodynamic Assessment of Nb-Ti System”,

Calphad, 25, 305-317 (2001) (Equi. Diagram, Thermodyn. Calculation, 42)

[2002Gho] Ghosh, G., “Thermodynamic and Kinetic Modeling of the Cr-Ti-V System”, J. Phase

Equilib., 23(4), 310-328 (2002) (Thermodyn., Equi. Diagram, 110)

[2002Tho] Thoma, D.J., Nibur, K.A., Chen, K.C., Cooley, J.C., Dauelberg, L.B., Hults, W.L., Kotula,

P.G., “The Effect of Alloying on the Properties of (Nb,Ti)Cr2 C15 Laves Phases”, Mater.

Sci. Eng. A, 329-331, 408-415 (2002) (Crys. Structure, Experimental, Phys. Prop., Equi.

Diagram, #, *, 25)

[2004Iva1] Ivanchenko, V, “Cr-Nb (Chromium-Niobium)”, MSIT Binary Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH,

Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 32)

[2004Iva2] Ivanchenko, V., “Cr-Ti (Chromium-Titanium)”, MSIT Binary Evaluation Program, in


GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 22)

[2004Zha] Zhao, J.C., Jackson, M.R., Peluso, L.A., “Mapping of the Nb-Cr-Ti Phase Diagram Using

Diffusion Multiples”, Z Metallkd., 95, 142-146 (2004) (Experimental,

Equi. Diagram, *, #, 34)

190


MSIT®

Cr–Nb–Ti


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

, (Cr,Nb, Ti)

(Cr)

1863

(Nb)

1455

( Ti)(h)

1670 - 882

cI2

Im3m

W a = 288.4

a = 330.04

a = 330.65

pure Cr at 27°C [V-C2]

pure Nb at 25°C [V-C2]

pure Ti [Mas2]

( Ti)(r)

882

hP2

P63/mmc

Mg

a = 295.06

c = 468.25


Nb1-xTixCr2(h)

NbCr2(h)

1730 - 1585

TiCr2(h)

1270 - 800

hP12

P63/mmc

MgZn2 a = 493.1

c = 812.3

a = 493.1

c = 800.5

C36 Laves phase.

at 66.7 at.% Cr and 25°C [2004Iva1];

solid solubility 62.2 to 70 at.% Cr.

at 25°C [V-C2].

Nb1-xTixCr2(r)

NbCr2(r)

1625

TiCr2(r)

1220

cF24

Fd3m

MgCu2 a = 693.88

a = 698.82

a = 699.49 to702.25

a = 693.2

0 x 1; C15 Laves phase.

at Nb5Ti28Cr67 and 25°C [1998Che]

at Nb20.1Ti17.3Cr62.6 and 25°C [2002Tho]

solid solubility from 61 to 69 at.% Cr

[2004Iva1].

at TiCr1.9 and 25°C [V-C2].

TiCr2(h)

1370 - 1270

hP24

P63/mmc

MgNi2

a = 493.2 0.2

c = 1601.0 0.1

C14 Laves phase.

at Ti1.12Cr2 and 25°C [V-C2].

191


MSIT®

Cr–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb

Cr Data / Grid: at.%

Axes: at.%

(Nb)

L

L+(Nb)

Fig. 1: Cr-Nb-Ti.


1900°C

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

(Nb)

L+(Nb)

L

L+(Cr)

(Cr)Fig. 2: Cr-Nb-Ti.


1800°C

192


MSIT®

Cr–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%(Cr)

L

L+β+βNbCr2(h)

β

βNbCr2(h)

L+(Cr)

L+β

(Nb)

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%(Cr)

L+(Cr)+βNbCr2(h)

βNbCr2(h)

αNbCr2(r)

L+β+αNbCr2(r)L

(βTi) (Nb)

β

Fig. 3: Cr-Nb-Ti.


1700°C

Fig. 4: Cr-Nb-Ti.


1600°C

193


MSIT®

Cr–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

β

(βTi) (Nb)

(Cr)

αNbCr2(r)

βTiCr2(h)

β+βTiCr2(h)+αNbCr

2(r)


2(r)

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

β

(Nb)


2(r)

L+β+βTiCr2(h)

L+(Cr)+βTiCr2(h)

βTiCr2(h)

(βTi)

αNbCr2(r)

L

(Cr)


2(r)

Fig. 6: Cr-Nb-Ti.


1400°C

Fig. 5: Cr-Nb-Ti.


1500°C

194


MSIT®

Cr–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

(Nb)(βTi)

β+βNbCr2(h)+αNbCr

2(r)

αNbCr2(r)

β

γTiCr2(h)

βTiCr2(h)

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

(βTi) (Nb)

(Cr)+αNbCr2(r)

αNbCr2(r)

βTiCr2(h)

(Cr)

β

(Cr)+βTiCr2(h)+αNbCr

2(r)


2(r)

Fig. 7: Cr-Nb-Ti.


1300°C

Fig. 8: Cr-Nb-Ti.


1250°C

195


MSIT®

Cr–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

β

(βTi) (Nb)

(Cr)

αNb1-x

TixCr

2(r)

βTiCr2(h)

(Cr)+αNb1-x

TixCr

2(r)+βTiCr

2(h)

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%(Cr)

βTiCr2(h)

αNb1-x

TixCr

2(r)

β

(βTi) (Nb)

Fig. 9: Cr-Nb-Ti.


1200°C

Fig. 10: Cr-Nb-Ti.


1150°C

196


MSIT®

Cr–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

(Nb)(βTi)

β

(Cr)

αNb1-x

TixCr

2(r)

βTiCr2(h)

(Cr)+βTiCr2(h)+αNb

1-xTi

xCr

2(r)

Fig. 12: Cr-Nb-Ti.


950°C

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%(Cr)

αNb1-x

TixCr

2(r)

βTiCr2(h)

β

(βTi) (Nb)

(Cr)+αNb1-x

TixCr

2(r)+βTiCr

2(h)

Fig. 11: Cr-Nb-Ti.


1000°C

197


MSIT®

Cr–Nb–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

(Nb)

β

(αTi)

(αTi)+β

αNb1-x

TixCr

2(r)

(Cr)

20

40

60

80

20 40 60 80

20

40

60

80

Ti Nb


Axes: at.%

αNb1-x

TixCr

2(r)

(Nb)

(Nb)+αNb1-x

TixCr

2(r)

(αTi)+αNb1-x

TixCr

2(r)+(Nb)

(αTi)+αNb1-x

TixCr

2

(αTi)+(Nb)(αTi)

(Cr)

Fig. 13: Cr-Nb-Ti.


800°C

Fig. 14: Cr-Nb-Ti.

Partial isothermal

section at 600°C

198


MSIT®

Cr–Nb–Ti

60 40 20

500

750

1000

1250

1500

1750

Ti 0.00

Nb 33.30

Cr 66.70

Ti

Cr, at.%

Te

mp

era

ture

, °C

L+(βTi)+βNbCr2(h)

(βTi)+βNbCr2(h)

(βTi)+βNbCr2(h)+αNbCr2(r)

(βTi)+αNbCr2(r)

(βTi)

(αTi)

βNbCr2(h)+αNbCr2(r)

L+(βTi)

(αTi)+(βTi)+αNbCr2(r)

(αTi)+(βTi)

αNbCr2(r)

βNbCr2(h)

L+βNbCr2(h)

(αTi)+αNbCr2(r)

LFig. 15: Cr-Nb-Ti.

Polythermal section

Ti-NbCr2

20 40 60

750

1000

1250

1500

1750

Ti 88.59

Nb 11.41

Cr 0.00

Ti 26.20

Nb 3.38

Cr 70.42Cr, at.%

Te

mp

era

ture

, °C

L

(αTi)+(βTi,Nb)

αNb1-xTixCr2(r)

(αTi)+(βTi,Nb)+αNb1-xTixCr2(r) (αTi)+αNb1-xTixCr2(r)

(βTi,Nb)

(βTi,Nb)+βTiCr2(h)

βTiCr2(h)+αNb1-xTixCr2(r)

(βTi,Nb)+βTiCr2(h)+αNb1-xTixCr2(r)

L+(βTi,Nb)

βTiCr2(h)

L+(βTi,Nb)+βTiCr2(h)

L+βTiCr2(h)

(βTi,Nb)+αNb1-xTixCr2(r)

Fig. 16: Cr-Nb-Ti.

An isopleth at a


Nb:Ti=1:4

199


MSIT®

Cr–Nb–Ti

10 20 30 40

500

750

1000

1250

1500

1750

Ti 75.00

Nb 25.00

Cr 0.00

Ti 37.50

Nb 12.50

Cr 50.00Cr, at.%

Te

mp

era

ture

, °C

L


(βTi,Nb)

Fig. 17: Cr-Nb-Ti.

An isopleth at a

constant atomic ratio

of Nb:Ti=1:3

20 40 60

750

1000

1250

1500

1750

Ti 74.43

Nb 25.57

Cr 0.00

Ti 20.30

Nb 6.98

Cr 72.72Cr, at.%

Te

mp

era

ture

, °C

(αTi)+(βTi,Nb)+αNb 1-xTixCr2(r)

(Ti,Nb)+αNb1-xTixCr2(r)

L

L+(βTi,Nb)

(βTi,Nb)+

(αTi)+(βTi,Nb)

(αTi)+αNb1-xTixCr2(r)

βTiCr2(h)

L+(βTi,Nb)+βTiCr2(h)

αNb1-xTixCr2(r)+βTiCr2(h)

αNb1-xTixCr2(r)+

L+βTiCr2(h)

(βTi,Nb)

αNb1-xTixCr2(r)

βTiCr2(h)

Fig. 18: Cr-Nb-Ti.

An isopleth at a


Nb:Ti=2:3

200


MSIT®

Cr–Nb–Ti

10 20 30 40

500

750

1000

1250

1500

1750

Ti 50.00

Nb 50.00

Cr 0.00

Ti 25.00

Nb 25.00

Cr 50.00Cr, at.%

Te

mp

era

ture

, °C

(βTi,Nb)


Fig. 19: Cr-Nb-Ti.

An isopleth at a


of Nb:Ti=1:1

10 20 30 40

500

750

1000

1250

1500

1750

Ti 25.00

Nb 75.00

Cr 0.00

Ti 12.50

Nb 37.50

Cr 50.00Cr, at.%

Te

mp

era

ture

, °C

(βTi,Nb)


Fig. 20: Cr-Nb-Ti.

An isopleth at a


of Nb:Ti=3:1

201


MSIT®

Cr–Nb–Ti

80 60 40

1600

1700

1800

Cr 94.59

Nb 0.00

Ti 5.41

Cr 23.01

Nb 68.66

Ti 8.33Cr, at.%

Te

mp

era

ture

, °C

L

(βTi,Cr)

βNbCr2(h)

(Nb)+βNbCr2(h)

L+(Nb)

(Nb)

L+(βTi,Cr)

L+βNbCr2(h)

L+(Nb)+βNbCr2(h)

(βTi,Cr)+βNbCr2(h)

L+(βTi,Cr)+βNbCr2(h)

Fig. 21: Cr-Nb-Ti.

An isopleth at a

constant Ti content of

5 mass%

202


MSIT®

Cr–Ni–Ti

Chromium – Nickel – Titanium

Nathalie Lebrun

Literature Data

No ternary phase has been found in the Cr-Ni-Ti system. Several isothermal sections have been investigated

experimentally using X-ray diffraction, micrographic analysis, triple diffusion technique and EPMA.

[1951Tay] determined the partial equilibrium diagrams in the Ni-rich part at 750, 1000 and 1150°C.

Additional isothermal sections at 850 and 927°C were reported by [1997Xu, 1998Bee].

All the binary phases have extension into the ternary region.

A wide (Ni) region was observed [1951Tay, 1997Xu, 1998Bee] and the solid solution widens as the

temperature increases [1951Tay]. The solid extension was found to be 10 at.% Cr at 1000°C [1955Kor1]

using a lattice parameter method on ternary alloys containing 20 at.% of Cr. This result agrees with those

of [1951Tay] and was confirmed later by [1997Xu]. The phase boundaries at 800°C for alloys containing

10 and 20 mass% Cr, determined by [1956Kor], also agree well with the results of [1951Tay] and are in

agreement with those of [1998Bee].

The ( Ti) phase region extended from the Cr-Ti binary to about 8-10 at.% Ni. The ( Cr) phase has a

homogeneity range which extends to about 8 at.% Ti from the Cr-Ni binary [1997Xu, 1998Bee].

The solid state solubility of Cr in TiNi3 was determined to be up to 7 at.% Cr [1998Bee]. The homogeneity

range of TiNi2 was extended up to 9 at.% Cr at 850°C [1998Bee] and 10 at.% at 927°C [1997Xu].

[1998Bee] confirmed the existence of a solubility range for TiNi (up to 9 at.% Cr) suggested previously by

[1997Xu].

The solubility of Ni in TiCr2 is about 10 at.% at 850°C whereas that for TiCr2 does not exceed a value

of 4 at.% [1998Bee]. [1997Xu] considered the existence of one phase for TiCr2 and found a solubility of Ni

in TiCr2 of about 1.5 at.% Ni.

Isothermal sections at 1027, 1277 and 1352°C were derived from thermodynamic calculations of phase

equilibria [1974Kau]. The large extension of the terminal phases (Ni), ( Cr) and ( Ti) was confirmed.

Comparison with the experimental data available [1951Tay, 1997Xu, 1998Bee] gave good agreement.

From the basis of the isothermal section calculated at 1027°C [1974Kau] and the partial one measured by

[1997Xu] at 927°C, [2003Gup] suggested in its assessment the phase equilibria which could exist at 927°C.

The schematic isothermal section is in good agreement with the ones reported by [1998Bee] in the Ni

corner, whereas discrepancies have been found concerning the phase equilibria involving ( Ti), Ti2Ni,

TiNi, TiCr2 and TiCr2.

From the binary systems and partial isothermal sections from [1951Tay, 1956Tay], [1986Gup] suggested a

schematic liquidus projection, except in the Ti-rich corner since no experimental data are available. Using

EPMA, DTA and metallographic techniques, [1978Hao] suggested the presence of a eutectic structure at

1220 2°C in the alloy of composition 35.4Ti-3.7Cr-60.9Ni (at.%).

A partial isopleth at 20 mass% Cr has been established by [1955Kor2] using thermal analysis and lattice

parameter method. Results are in general good agreement with the limit of the (Ni) solubility ranges

estimated by [1951Tay, 1974Kau, 1998Bee, 1997Xu]. Slight discrepancies are observed concerning the

L+(Ni) phase region between the experimental results of [1955Kor2] and the calculated isothermal section

of [1974Kau] at 1277 and 1352°C. [1988Nar] also determined a polythermal section at constant Ni content

of 8 mass% and Cr varying from 0 to 10 mass%. Good agreement of the ( Ti) limit solubility is observed

with experimental isothermal section established at 850°C by [1998Bee].

Binary Systems

A complete assessment of the binary Cr-Ti system was done by [1987Mur1]. It was reported an unknown

reaction at around 1270°C involving the TiCr2 and TiCr2 phases. Recently, [2000Zhu] calculated the

phase diagram using a thermodynamic description which reproduce well the experimental data from the

literature. The unknown reaction mentioned in [1987Mur1] was found by [2000Zhu] to be a peritectoid

203


MSIT®

Cr–Ni–Ti

( Ti)+ TiCr2 TiCr2 at 1271°C and a eutectoid TiCr2 TiCr2+( Cr) at 1269°C. The new reactions

proposed by [2000Zhu] are accepted in this assessment and the binary phase diagram is then a compilation of

the phase diagrams proposed by [1987Mur1] and [2000Zhu].

The binary Ni-Ti system has been extensively reviewed by [1991Mur]. More recently, [1996Bel] has done

a new assessment of the thermodynamic properties of the stable phases, based on thermochemical and phase

diagram data from the literature. Their calculation is in good agreement with the phase equilibria reported

by [1991Mur]. The solid homogeneity range of TiNi3 has been reproduced by [1996Bel] and is in good

agreement with the literature data. Using symmetric two sublattice model, [1999Tan] described the

transformation from the ordered TiNi phase to the disordered TiNi bcc phase which occurs at around

93.5°C. The accepted diagram in this assessment is then a compilation of [1991Mur] and [1996Bel].

The binary Cr-Ni system has been assessed by [1991Nas]. Later, [1995Tom] calculated the phase diagram

from its optimized thermodynamic data determined from Knudsen cell mass spectrometry measurements.

A slight difference is observed on the composition of the ( Cr) phase at the eutectic temperature 1345°C.

[1995Tom] reported a value of 64.73 at.% Cr instead of 69 at.% of Cr assessed by [1991Nas]. The value

reported by [1995Tom] is retained is this assessment. The binary diagram accepted in this assessment is

mainly based on the work of [1995Tom] with some modification at low temperature taking into account the

existence of the CrNi2 phase proposed by [1991Nas].

Solid Phases

Crystallographic data for all the solid phases are presented in Table 1. [1992Shi] studied various ternary

alloys with the composition of Ti50-xCrxNi50, Ti50-x/2CrxNi50-x/2 and Ti50CrxNi50-x (0 < x 3) annealed 24

hours at 1000°C. In the two first alloys, Cr atoms occupy both the sites Ni and Ti with nearly equal fractions,

whereas the Ni sites are preferentially occupied by Cr atoms in the third alloy.


A monovariant eutectoid transformation ( Ti) ( Ti)+Ti2Ni+ TiCr2 takes place at 650°C [1988Nar] (see

Table 2). [1978Hao] reported the existence of a eutectic at 1220°C with a composition of

35.38Ti-3.67Cr-60.95Ni (at.%). The phases involved in this eutectic have not been detected precisely.

Consequently, this result has not been retained in this assessment.

Liquidus Surface

From the binary systems and the isothermal sections reported by [1951Tay, 1956Tay, 1997Xu, 1974Kau,

1998Bee], a schematic liquidus surface has been suggested by [1990Gup]. This liquidus projection is

uncertain since no experimental data are available. Consequently, the schematic liquidus surface and the

reaction scheme proposed by [1990Gup] were not retained in this assessment.

Isothermal Sections

Slight modifications on the binary boundaries and inside the ternary region have been done on all the

isothermal sections accepted in this assessment in order to be in agreement with the accepted binary systems

and the composition-temperature sections accepted in this assessment (see Figs. 1a and 1b).

Two isothermal sections at 850 and 927°C have been determined by [1998Bee] and [1997Xu], respectively.

Only the isothermal section at 850°C have been accepted in this assessment and reported in Fig. 2. The not

well known single phase regions and position of the lines which delimited the two- and the three-phase

fields are indicated as dashed lines in the diagram. The isothermal section at 927°C has not been retained in

this assessment since no indications have been reported on the phase field equilibria.

The calculated isothermal sections at 1027, 1277 and 1352°C taken from [1974Kau] are shown in Figs. 3,

4 and 5, respectively. In agreement with the binary systems, the phase fields involving the TiCr2, TiCr2,

( Ti), ( Cr) and liquid phases have been modified and are shown as dashed lines in Fig. 3. Slight

modification of the L + (Ni) region to be in agreement with the experimental temperature-composition

204


MSIT®

Cr–Ni–Ti

section measured by [1955Kor2]. The partial isothermal sections published by [1951Tay] are in good

agreement with the ones reported in Figs. 2 and 3.


Two polythermal sections has been reported, at 8 mass% Ni by [1988Nar] (Fig. 1a) and at 20 mass% Cr by

[955Kor2] (Fig. 2). General agreement is observed between these experimental isopleths and the isothermal

sections measured by [1951Tay, 1997Xu, 1998Bee] and those calculated by [1974Kau]. However,

discrepancies are observed between the binary Ni-Ti the polythermal section reported on Fig. 1a. The curves

have been modified in agreement with the accepted Ni-Ti diagram. As strong discrepancies are observed

the corresponding curves have been indicated as dashed lines. [1955Kor2] measured a solid solubility of Ti

in (Ni) of about 7 mass% Ti and 20 mass% Cr. A three phase field L+(Ni)+TiNi3 was detected.


TiNi based alloys are of great importance because of superior corrosion and wear resistance. A shape

memory effect can cause uncontrollable shape recovery after machining, bending and rolling. To avoid this

problem, it is necessary to suppress martensitic transformation, i.e. to stabilize the stable TiNi phase with a

CsCl structure. [1998Hos] reported the effect of Cr addition on the martensitic and the austenite

transformations and mechanical properties of alloys with a constant at.% Ti and Cr and Ni varying from 0

to 50 at.%. The authors also studied the yield stress and hardening properties from 77 to 873 K for the TiNi

phase in alloys based on Ni-49 mol% Ti.

[1950Cra] also reported the effect of Ni on the mechanical properties (elongation, tensile strength, Vickers

hardness) of as-rolled Cr-Ti alloys. [1955Kor2] determined the high temperature strength of alloys Ti-20Cr

up to 10Ni (mass%).

[1988Nar] also investigated the corrosion resistance and the electrochemical behavior of alloys of the

Cr-Ni-Ti system in 10 % NaCl, NaOH and HCl solutions.

Miscellaneous

[1968Che] studied the activity of Ti in nickel melts containing 5, 18 and 25 % of Cr at 1600°C. The activity

coefficient of Ti in (Ni) is 0.0002 and varies linearly with Cr content.

Partitioning of Cr between (Ni) and TiNi3 phases was determined by the diffusion couple method between

1000 and 1200°C [1994Jia]. It was found that the partitioning coefficient varies from 0.22 at 1000°C to 0.45

at 1200°C.

References

[1950Cra] Craighead, C.M., Simmons, O.W., Eastwood, L.W., “Ternary Alloys of Titanium”, Trans.

Am. Inst. Min. Metall. Eng., 188, 514-538 (1950) (Experimental, Mechan. Prop., 1)

[1951Tay] Taylor, A., Floyd, R.W., J. Inst. Met., 80, 577-587 (1951) (Equi. Diagram, Experimental,


[1955Kor1] Kornilov, I.I., Snetkov, A.Ya., “X-Ray Investigation of Limited Solid Solutions of Nickel”

(in Russian), Izv. Akad. Nauk SSSR, Otdel. Tekh. Nauk, 7, 84-88 (1955) (Experimental, 12)

[1955Kor2] Kornilov, I.I., Kosmodemyansky, V.V., “Relationships Between Composition,

Temperature, and High-st Rength”, Izv. Akad. Nauk SSSR, Otdel. Tekh. Nauk, 2, 90 (1955)


[1956Kor] Kornilov, I.I., Pryakhina, L.I., “The Composition Elevated Temperature Strength Diagram

of the Ni-Cr-Ti”, Izv. Akad. Nauk SSSR, Otdel. Tekh. Nauk, 7, 103-110 (1956) (Equi.


[1956Tay] Taylor, A., “Constitution of Nickel-Rich Quaternary Alloys of the Ni-Cr-Ti-Al System”,

J. Met., 1356-1362 (1956) (Equi. Diagram, Expermental, 8)

205


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Cr–Ni–Ti

[1968Che] Cherkasov, P.A., Averin, V.V., Samarin, A.M., “Activities of Silicon and Titanium in

Molten Iron, Cobalt, and Nickel Containing Chromium”, Russ. J. Phys. Chem., 42(3),

401-404 (1968) translated from Zh. Fiz. Khim, 42(3), 767 (1968) (Experimental, 19)

[1974Kau] Kaufman, L., Nesor, H., “Calculation of Superalloy Phase Diagrams: Part I”, Metall. Trans.,

5(7), 1617-1621 (1974) (Calculation, Equi. Diagram, 19)

[1978Hao] Haour, G., Mollard, F., Lux, B., Wright, I.G., “New Eutectics Based on Fe, Co and Ni”,

Z. Metallkd., 69, 149-154 (1978) (Experimental, 14)

[1986Gup] Gupta, K.P., Rajendraprasad, S.B., Jena, A.K., “The Chromium - Nickel - Titanium

System”, J. Alloy Phase Diagrams, 2(1), 31-37 (1986) (Review, 11)

[1987Mur1] Murray, J.L., “The Cr-Ti (Chromium-Titanium) System” in “Phase Diagrams of Binary

Titanium System”, ASM Internal, Metals Park, OH, 68-78 (1987) (Equi. Diagram, Crys.

Structure, Thermodyn., Review, 60)

[1987Mur2] Murray, J.L., “The Ni-Ti (Nickel-Titanium) System” in “Phase Diagrams of Binary



[1988Nar] Nartova, T.T., Mogutova, T.V., Volkova, M.A., Mikaberidze, M.P., Lordkipanidze, I.N.,

“Phase Equilibria and Corrosion Resistance of Ti-Ni-Cr Alloys”, Izv. Akad. Nauk SSSR, 3,

182-184 (1988) (Experimental, Mechan. Prop., Equi. Diagram, Crys. Structure, Electr.

Prop., 5)

[1990Gup] Gupta, K.P., “Ternary Systems Containing Chromium-Nickel, Copper-Nikel and

Iron-Nickel”, in “Phase Diagram of Ternary Nickel Alloys”, Indian Institute of Technology,

Calcutta, 93-102 (1990) (Review, 12)

[1991Mur] Murray, J.L., “The Ni-Ti (Nickel-Titanium) System”, in “Phase Diagrams of Binary

Titanum System”, ASM Internal, Metals Park, OH, 197-210 (191) (Equi. Diagram, Crys.

Struct., Thermodyn., Reveiew, 110)

[1991Nas] Nash, P., “Cr-Ni (Chromium-Nickel)” in “Phase Diagrams of Binary Nickel Alloys”, ASM

Internal, Metals Park, OH, 75-84 (1991) (Equi. Diagram, Crys. Structure, Thermodyn.,

Review, 126)


Martensitic Transformations”, Mater. Trans., JIM, 33(3), 165-177 (1992) (Calculation,

Crys. Structure, Theory, 101)

[1993Yan] Yan, Z.H., Klassen, T., Michaelsen, C., Oehring, M., Bormann, R., “Inverse Melting in the

Ti-Cr System”, Phys. Rev. B, 47(14), 8520-8529 (1993) (Equi. Diagram, Experimental, 30)

[1994Jia] Jia, C.C., Ishida, K., Nishizawa, T., “Partitioning of Alloying Elements Between (A1) and

(DO24) Phases in the Ni-Ti Base Systems”, in “Exp. Methods Phase Diagram Determ.”,

Morral, J.E., Schiffman, R.S., Merchant, S.M., (Eds.), The Minerals, Metals & Materials

Society, 31-38 (1994) (Equi. Diagram, Experimental, 8)

[1995Tom] Tomiska, J., Kopecky, K., Belegratis, M.S., Myers, C., “Thermodynamic Mixing Functions

and Phase Equilibria in the Nickel-Chromium System by High-Temperature Knudsen Cell

Mass Spectrometry”, Metall. Mater. Trans. A, 26A(2), 259-265 (1995) (Equi. Diagram,

Experimental, Thermodyn., 49)

[1996Bel] Bellen, P., Kumar, K.C.H., Wollants, P., “Thermodynamic Assessment of the Ni-Ti Phase

Diagram”, Z. Metallkd., 87(12), 972-978 (1996) (Review, Equi. Diagram, Thermodyn., 43)

[1997Xu] Xu, H.H., Jin, Z.P., “The Determination of the Isothermal Section at 1200 K of the Cr-Ni-Ti

Phase Diagram”, Scr. Mater., 37(2), 147-150 (1997) (Equi. Diagram, Experimental, 6)

[1998Bee] Beek, J.A., Kodentsov, A.A., Loo, F.J.J., “Phase Equilibria in the Ni-Cr-Ti System at

850°C”, J. Alloys Compd., 270, 218-223 (1998) (Equi. Diagram, Experimental, 9)

[1998Hos] Hosoda, H., Hanada, S., Inoue, K., Fukui, T., Mishima Y., Suzuki, T., “Martensite

Transformation Temperatures and Mechanical Properties of Ternary NiTi Aalloys with

Offstoichiometric Compositions”, Intermetallics, 6(4), 291-301 (1998) (Experimental,

Mechan. Prop., 24)

206


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[1999Tan] Tang, W., Sundman, B., Sandstroem, R., Qiu, C., “New Modelling of the B2 Phase and its

Associated Martensitic Transformation in the Ni-Ti System”, Acta Mater., 47(12),

3457-3468 (1999) (Thermodyn., Calculation, 51)

[2000Zhu] Zhuang, W., Shen, J., Liu, Y., Shang, L., Du, Y., Schuster, J.C., “Thermodynamic

Optimization of the Cr-Ti System”, Z. Metallkd., 91, 121-127 (2000) (Thermodyn.,

Calculation, 53)

[2003Gup] Gupta, K.P., “The Cr-Ni-Ti (Chromium-Nickel-Titanium) System-Update”, J. Phase

Equilib., 24(1), 86-89 (2003) (Equi. Diagram, Review, 7)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Cr,Ti)

< 1863

( Cr)

< 1863

( Ti)

1670 - 882

cI2

Im3m

W

a = 328 - 3.98xcr

a = 288.48

a = 330.65

0 to 100 at.% Cr at 25°C quenched solid

solution [1993Yan]

pure Cr at 25°C [Mas2]


( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

pure Ti at 25°C [Mas2] dissolves 0.6

at.% Cr at 667°C [1987Mur1]

(Ni)

< 1455

cF4

Fm3m

Cu a = 352.40

86.1 to 100 at.% Ni [1991Nas]

pure Ni at 25°C [Mas2]

TiCr2 (h2)

1359 - 1269

hP12

P63/mmc

MgZn2 a = 493.2 0.2

c = 800.5 0.3

63.90 to 65.7 at.% Cr at 1271°C

[2000Zhu]

Ti1.12Cr2 [V-C2]

TiCr2 (h1)

1271 - 804

hP24

P63/mmc

MgNi2

a = 493.2 0.2

c = 1601 0.1

63.8 to 66 at.% Cr at 1223°C [2000Zhu]

Ti1.12Cr2 [V-C2]

TiCr2 (r)

< 1223

cF24

Fm3m

Cu2Mg

a = 693.2 0.4 62.8 to 66.5 at 804°C [2000Zhu]

TiCr1.9 [V-C2]

TiNi3< 1380

hP16

P63/mmc

Ni3Ti

a = 510.28

c = 827.19

75 to 80.1 at.% Ni at 1300°C [1996Bel]

[V-C2]

TiNi

< 1310

cP2

Pm3m

CsCl

a = 301.0

49.5 to 57 at.% Ni [1987Mur2]

Ti0.98Ni1.02 [V-C2]

Ti2Ni

< 984

cF96

Fd3m

NiTi2

a = 1132.4

33 to 34 at.% Ni [1987Mur2]

Annealed at 950°C for 72 hours [V-C2]

CrNi2< 590

oI6

MoPt2(?) a = 252.4

b = 757.1

c = 356.8

23.5 to 40 at.% Cr at 500°C [1991Nas]

[P] (c = a0 of solid solution of (Cr,Ni)

subcell at 66.3 at.% Ni)

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Cr–Ni–Ti



Cr Ni Ti

( Ti) ( Ti) + Ti2Ni + TiCr2 650 E ( Ti)

( Ti)

Ti2Ni

TiCr2

0.0

0.0

0.0

66.7

0.0

0.0

33.3

0.0

100.0

100.0

66.7

33.3

90

500

600

700

800

900

1000

Ti 93.38

Cr 0.00

Ni 6.62

Ti 83.90

Cr 9.42

Ni 6.68Ti, at.%

Te

mp

era

ture

, °C

(βTi)

(βTi) + Ti2Ni

(αTi)+Ti2Ni

(αTi) + (βTi) + Ti2Ni

(αTi) + Ti2Ni + αTiCr2

Fig. 1a: Cr-Ni-Ti.

Isopleth at constant

Ni content of 8

mass%

208


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Cr–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Cr


Axes: at.%

(Ni)

(βCr)

TiNi3

TiNi

Ti2Ni

(βTi)

(αTi) αTiCr2

βTiCr2

Fig. 2: Cr-Ni-Ti.


850°C

10

600

700

800

900

1000

1100

1200

1300

1400

1500

Ti 0.00

Cr 22.01

Ni 77.99

Ti 11.69

Cr 21.54

Ni 66.77Ti, at.%

Te

mp

era

ture

, °C

(βTi)

L

L + TiNi3

L + (βTi) + TiNi3L + (βTi)

Fig. 1b: Cr-Ni-Ti.


mass% Cr

209


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Cr–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Cr


Axes: at.%

(Ni)

(βCr)

TiNi3

TiNi

L

(βTi)

αTiCr2

βTiCr2

Fig. 3: Cr-Ni-Ti.


1027°C

20

40

60

80

20 40 60 80

20

40

60

80

Ti Cr


Axes: at.%

(βCr)

(Ni)

TiNi3

NiTi L

(βTi)

γTiCr2

Fig. 4: Cr-Ni-Ti.


1277°C

210


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Cr–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Cr


Axes: at.%

(Ni)

(βCr)(βTi)

TiNi3

L

γTiCr2

Fig. 5: Cr-Ni-Ti.


1352°C

211


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Cu–Mg–Ni

Copper – Magnesium – Nickel

Hans Leo Lukas, Lazar Rokhlin

Literature Data

Many research groups dealt with the constitution of the Cu-Mg-Ni system [1951Koe, 1952Lie, 1956Mik1,

1956Mik2, 1972Feh, 1972Kom1, 1983Dar, 1983Kar, 1986She, 1995Ips]. Reviews were published by

[1939Vos, 1949Jae, 1977Ray, 1979Cha, 1979Dri, 1995Ips]. Early work in this ternary system was based

on the assumption of continuous solid solutions among the corresponding pairs of binary compounds

[1939Vos, 1951Koe, 1956Mik1, 1956Mik2]. A major breakthrough was achieved when the section

Cu2Mg-Ni2Mg was recognized as a pseudobinary peritectic system [1949Jae, 1952Lie, 1972Feh, 1983Dar,

1983Kar, 1995Ips]. Detailed crystallographic inspection established a ternary compound

(Ni0.45Cu0.55)2Mg as a new stacking variant of the Laves phases [1972Kom1, 1972Kom2, 1974Kri]. There

is, however, still some lack of information on the complete incorporation of this compound into the phase

diagram.

[1949Jae] reported solid solubilities of about 15 mol% Cu2Mg in Ni2Mg and about 8 mol% Ni2Mg in

Cu2Mg, but later works considered these data not as reliable. [1951Koe] investigated the phase relations in

three isopleths for Cu:Ni = 3:1, 1:1 and 1:3 by thermal analysis of 40 alloys. He used the results for a basic

construction of the liquidus surface, which shows three monovariant troughs, directed from the Mg-Ni to

the Cu-Mg side, so that in pairs the binary eutectic transformations L Ni+Ni2Mg with L Cu+Cu2Mg, the

peritectic transformation L+Ni2Mg NiMg2 with the eutectic transformation L Cu2Mg+CuMg2, and the

eutectic transformations L NiMg2+(Mg) and L CuMg2+(Mg) continuously turn one into the other

without any invariant reaction. He was aware, that this form of the liquidus surface is a simplification,

neglecting the non-isomorphous lattices of the corresponding pairs of compounds Ni2Mg-Cu2Mg and

NiMg2-CuMg2.

[1952Lie] used thermal analysis, microscopic and X-ray methods for a construction of the polythermal

section Cu2Mg-Ni2Mg, which was shown to be a pseudobinary peritectic system: L+Ni2Mg Cu2Mg. The

solubility limits of the Ni2Mg- and Cu2Mg- based solid solutions were determined. [1956Mik1] constructed

the liquidus surface of the Cu-Mg-Ni phase diagram from thermal analysis and microscopic observations.

These authors, like [1951Koe] treated the two phases Ni2Mg and Cu2Mg as a continuous solid solution. In

the Mg-rich part of the system [1956Mik1] reported two invariant four-phase reactions,

L+(Ni,Cu)2Mg NiMg2+CuMg2 at 540°C and L (Mg)+NiMg2+CuMg2 at 480°C. For the first one the

composition of liquid is given as Cu33.5Mg65Ni1.5, however, this is incompatible with Raoult’s law, which

gives an initial slope of the liquidus of about 5 K/at.% Ni starting at the binary CuMg2 compound,

congruently melting at 568°C, whereas the values given by [1956Mik1] correspond to 18 K/at.% Ni. As

furthermore CuMg2 dissolves about 1 at.% NI, this slope should be related to the difference of the Ni

contents of liquid and CuMg2 and thus the temperatures of congruent melting of CuMg2 and a liquidus point

of the L+CuMg2 equilibrium at this Ni content must be even closer. [1956Mik1] furthermore supposed a

ternary compound NiCuMg, based on electric resistivity measurements on alloys of the Cu2Mg-Ni2Mg

section by [1956Mik2]. This compound may be identified with the stacking variant of the Laves phases

found by [1972Kom1, 1972Kom2].

[1972Feh] investigated the Cu corner of the Cu-Mg-Ni phase diagram along the monovariant eutectic line

starting at the binary eutectic L (Cu)+Cu2Mg using thermal analysis, electron microprobe analysis,

microscopic and X-ray methods. They established the invariant four-phase reaction

L+Ni2Mg (Cu,Ni)+Cu2Mg. [1972Feh], like [1952Lie], considered the Cu2Mg-Ni2Mg section to be a

pseudobinary system of two solid phases Cu2Mg and Ni2Mg. However, they assumed the solubility of

Cu2Mg in the Ni2Mg phase to decrease rapidly with decreasing temperatures. About 4 to 7 at.% Cu were

measured by microprobe analyses of this phase in alloys homogenized 70 h at 700°C. [1972Feh] tentatively

outlined a reaction scheme taking into account the liquidus temperatures and invariant reactions reported by

[1956Mik1]. They also constructed an isothermal section at 475°C revealing the phase equilibria in solid

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state, however the later detected large solubility of CuMg2 in NiMg2 was not yet considered. In addition

[1972Feh] constructed four partial isothermal sections at 850, 808, 800, and 750°C from the Cu-Ni side up

to the Ni2Mg-Cu2Mg line. These sections must be taken as sketches to explain the four-phase reaction,

rather than as quantitative diagrams, especially regarding the Mg solubility in the (Ni,Cu) solid solution.

The limits at the Cu-Mg binary are about 0.5 at.% larger than in the accepted binary Cu-Mg phase diagram.

For digitizing a small figure this may be taken as good agreement, but the limiting solubility of Mg in

(Ni,Cu) at the (Ni,Cu)+Ni2Mg two-phase field at 40 at.% Ni is drawn to increase with decreasing

temperature from about 4 at.% Mg at 850°C to 5 at.% Mg at 730°C. This is not likely and no evidence of

experimental support for these values is given in the paper. For the Mg-rich part of the system with more

than 33 at.% Mg [1972Feh] constructed a tentative reaction scheme accepting the four-phase equilibria

published by [1956Mik1].

[1972Kom1] investigated details of the crystal structure of Ni2Mg-Cu2Mg alloys within the range 50 to 55

mol% Cu2Mg by single crystal X-ray photographs. The alloys were annealed at temperatures between 500

and 800°C. The authors revealed a hexagonal ternary phase at 55 mol% Cu2Mg: (Ni0.45Cu0.55)2Mg as a new

stacking variant of the Laves phase structures. [1972Kom2, 1974Kri] explained the formation of this ternary

phase as a function of the electron concentration. The conclusions of [1972Kom1, 1972Kom2] eventually

correspond to the suggestions of [1956Mik1, 1956Mik2] about the compound “NiCuMg”.

The reviews by [1977Ray, 1979Cha, 1979Dri] essentially accepted the limiting solubilities of the Laves

phases from [1972Feh] and rejected those of [1952Lie].

[1983Dar] investigated alloys along the line NiMg2-CuMg2 by X-ray powder diffraction and established

the formation of an extended NiMg2-based solid solution (up to 85 mol% CuMg2 at 600°C) with linear

variation of the unit cell parameters. At higher Cu concentrations the NiMg2 solid solution coexists with

practically pure CuMg2.

[1983Kar] used microscopic and X-ray analyses for the construction of an isothermal section at 400°C in

the 40-100 mass% Mg area. The investigation was based on a number of prepared alloys which showed only

the phases NiMg2 and CuMg2 in equilibrium with the Mg solid solution. No measurable solubility of Cu

and Ni in solid magnesium was found. The solubility of Cu in solid NiMg2 along the NiMg2-CuMg2 line

was reported to be quite high whilst the solubility of Ni in solid CuMg2 was reported to be quite small. These

conclusions of [1983Kar] agree with the results of [1983Dar]. The extensions of the NiMg2 and CuMg2

homogeneity areas across the NiMg2-CuMg2 line were reported by [1983Kar] to be rather narrow (at least

less than 1 at.%).

[1986She] prepared the ternary alloy Ni0.75Cu0.25Mg2 by chemical reaction at 560-580°C without fusion

resulting in a dark grey powder. X-ray powder diffraction proved solid solution of Cu in NiMg2. This result

confirms once more the high solubility of Cu in NiMg2.

[1995Ips] reinvestigated experimentally the whole Cu-Mg-Ni phase diagram employing differential

thermal analysis, X-ray powder diffraction and isopiestic vapor pressure measurements. Four polythermal

sections were constructed: isopleths with constant xCu/xNi ratios of 2.0, 1.0 and 0.5 and at constant

magnesium content of 71 at.%. [1995Ips] confirmed the invariant reactions U1, U2 and E1 reported by

[1972Feh] and accepted U3. They assessed a table giving temperatures and compositions of the phases

participating in all invariant four-phase reactions. These data, however, disagree to some extent with the

liquidus surface constructed by [1956Mik1], especially in the Mg corner.

Thermodynamic investigations were performed by [1991Gna, 1993Gna1, 1993Gna2, 1994Gna, 1995Feu,

1995Ips].

Thermodynamic assessments were reported by [1995Feu, 1995Jac, 2002Gor]. The first two are restricted

to modeling of the liquid phase, ignoring the ternary solubilities in the solid phases. [2002Gor] reported a

complete ternary dataset, but there seem to be errors in the reported values. An attempt to reproduce the

published calculated diagrams by these data resulted in significantly deviating diagrams.

Three papers [1995Cho, 1996Gon, 1997Gan] constructed formulas to predict ternary thermodynamic

properties from the binary ones and applied them to the Cu-Mg-Ni system.

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Cu–Mg–Ni

Binary Systems

The three binary systems Cu-Mg, Cu-Ni, and Mg-Ni are accepted from [Mas2]. Thermodynamic

assessments of the three binary systems were prepared in the COST 507 action [1998Ans]. The phase

diagrams calculated from these assessments agree very well with those of [Mas2].

Solid Phases

One ternary phase was established [1972Kom1, 1972Kom2] in the Ni2Mg-Cu2Mg section close to 50

mol%, but its range of stability was not fully determined, neither with respect to temperature nor to

composition. [1998Tsu] confirmed an alloy molten from equiatomic parts Cu+Mg+Ni to consist of this

phase.

(Cu) and (Ni) form a continuous solid solution. Three of the four binary phases form solid solutions of

substantial extensions along the sections Ni2Mg-Cu2Mg and NiMg2-CuMg2. The mutual solubility limits

of Ni2Mg and Cu2Mg are accepted from [1952Lie]. These data were obtained from lattice parameter

measurements by X-ray diffraction and may be considered as quite reliable. [1972Feh] reported only 5-7

mol% solubility of Cu2Mg in Ni2Mg, derived from microprobe analysis of Ni in this phase in three-phase

samples of compositions Cu48Mg17Ni35 and Cu39Mg15Ni46, annealed at 800°C and quenched. The binary

Laves phases exhibit slightly extended homogeneity ranges: 4.3 at.% for Cu2Mg and about 0.7 at.% Mg for

Ni2Mg. The width across the 33.3 at.% Mg line in the ternary system was not investigated.

The solubility of CuMg2 in NiMg2 was reported as 28 at.% Cu at 600°C [1983Dar], 24 at.% Cu at 400°C

[1983Kar] or 25 at.% Cu at 450°C [1995Ips]. These data agree fairly well, also with [1986She]. The

solubility of NiMg2 in CuMg2 is negligible, [1983Kar] estimated it to be 1 at.% Cu at 400°C, whereas

[1983Dar] did not reveal it at all. The widths of the homogeneity ranges of NiMg2 and CuMg2 across the

NiMg2-CuMg2 line are practically zero [Mas, 1983Kar]. The solubility of Cu and Ni in solid (Mg) is very

small. In the binary Cu-Mg system it is less than 0.013 at.% Cu, for Ni no value was reported. All solid

phases are listed in Table 1.


The section Ni2Mg-Cu2Mg is recognized as a pseudobinary system. It is shown in Fig. 1, which reproduces

in general the findings of [1952Lie]. The range, where the ternary Laves phase may be stable is indicated

as hatched area according to [1972Kom1]. Corrections were made to meet the melting points of Ni2Mg and

Cu2Mg reported for the Mg-Ni and Cu-Mg binary phase diagrams [Mas2]. The liquidus and solidus lines

as well as the existence of a peritectic in this pseudobinary system, constructed by [1952Lie], have to be

considered as quite reliable. They were not disputed and were supported by [1972Feh, 1977Ray]. For the

extension of the two-phase field Cu2Mg+Ni2Mg the rather precise X-ray data of [1952Lie] were preferred

over those of [1972Feh], who gave a solubility of Cu2Mg in Ni2Mg decreasing much more with decreasing

temperature.


There are four invariant four-phase equilibria in the system and most probably two maxima of three-phase

equilibria. Their temperatures and phase compositions are given in Table 2. The reactions U1 and U2 were

first reported by [1972Feh] and experimentally verified by [1995Ips]. The compositions of the Ni2Mg phase

in Table 2 are adjusted to the data of [1952Lie]. Reaction U3 was first reported by [1956Mik1] as

L+Cu2Mg NiMg2+CuMg2 at 540°C. This reaction implies a three-phase field L+NiMg2+CuMg2 going to

lower temperatures and the authors located it about 1 at.% Ni behind the binary melting point maximum of

CuMg2 at 568°C. This is a severe contradiction to Raoult’s law, which predicts for 1 at.% Ni about 5 K

freezing point depression, using the melting enthalpy of CuMg2 from the accepted binary system [1998Ans]

and assuming zero solubility of Ni in CuMg2. With some solubility of Ni in CuMg2 an even smaller

temperature difference is expected. Therefore here this reaction is taken from a tentative calculation

described below in section Thermodynamics as L+NiMg2 Cu2Mg+CuMg2 at 553°C with composition of

L near the binary eutectic e4. The three-phase field L+NiMg2+CuMg2 passes by a maximum e3 at about

214


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Cu–Mg–Ni

1 at.% Ni distance the binary CuMg2 phase and then goes to E1. [1972Feh, 1995Ips] did not investigate U3

and adopted it from [1956Mik1]. From their calculation [2002Gor] reported very similar phase

compositions as given in Table 2, but a temperature of 559°C. E1 was first reported by [1956Mik1] and

experimentally verified by [1972Feh, 1995Ips], except the composition of NiMg2, which is taken from

[1983Dar, 1983Kar]. The data on the maximum p2 are accepted from [1952Lie], those for the maximum e3

are taken from the calculation described below in section thermodynamics. The reaction scheme is

presented in Fig. 2. Figure 3 shows the projection of the invariant equilibrium planes together with the lines

of double saturation of liquidus and solidus, calculated from the data of Table 3.

Liquidus Surface

Figure 4 shows the liquidus surface, calculated from the dataset given in Table 3. At lower Mg contents it

deviates slightly from the best experimental data, but, as the experiments cover only some restricted areas,

it seems to be not possible to construct a better self-consistent diagram of the whole liquidus surface.

Isothermal Sections

Figure 5 shows the calculated isothermal section at 475°C. It differs from that constructed by [1972Feh] by

the concentrations of the solid phases, especially CuMg2 and NiMg2 where the data of [1983Dar, 1983Kar]

are taken into account. The solubilities of the Laves phases across the 33.3 at.% Mg line must be taken as

tentative. They are extrapolations from the binary assessments of these phases. The (Ni,Cu) corner of the

(Ni,Cu)+Cu2Mg+Ni2Mg field was drawn by [1972Feh] more near to Cu and with higher Mg content.


Figure 6 displays a vertical section of the phase diagram, constructed after [1972Feh]. It follows the eutectic

groove from the binary eutectic point L Cu+Cu2Mg to the counterpart L Ni+Ni2Mg. Figure 7 displays the

vertical section for the constant ratio Cu:Ni = 1:1 (at.%), and Fig. 8 displays the vertical section for a

constant Mg content of 71 at.%. The diagrams in Figs. 7 and 8 are calculated using Table 3. Figure 7 above

50 at.% Mg and Fig. 8 agree well with the experimental points of [1995Ips], The extension of the

three-phase field L+NiMg2+CuMg2 by [1995Ips] was drawn much smaller, but, by dashed lines the authors

themselves indicated that as tentative. Figure 7 below 50 at.% Mg shows somewhat higher temperatures

than [1995Ips] and there it has to be taken as tentative.

Thermodynamics

Thermodynamic properties of ternary Cu-Mg-Ni alloys were determined from isopiestic magnesium vapor

pressure measurements in the temperature range from 777 to 1077°C along three isopleths with

xCu/xNi = 2.0, 1.0 and 0.5 between about 20 and 90 at.% Mg. Thermodynamic activities and partial molar

Gibbs energies of magnesium were derived for the liquid phase and integral Gibbs energies of formation

were calculated by Gibbs-Duhem integration. The composition dependence of the activities is reported for

the three isopleths [1991Gna, 1993Gna1, 1993Gna2, 1995Ips].

[1972Pre] determined the enthalpy of formation of solid alloys along the section Cu2Mg-Ni2Mg within 0 to

40 mol% Ni2Mg. Behind a minimum at 10 mol% the enthalpy increases with increasing Ni2Mg content.

Enthalpies of liquid Cu-Mg-Ni alloys were studied by [1995Feu] using various types of calorimeters to

determine the integral enthalpies of mixing and heat capacities.

[1995Jac] performed a thermodynamic calculation of the ternary system and reported a partial diagram of

the isopleth at xCu/xNi = 0.5, compared with experimental points determined by [1995Ips]. These authors

used the thermodynamic datasets of the binary systems of the COST 507 action [1998Ans] and added a

ternary term to the Gibbs energy of liquid. They did not consider the ternary solubilities in the solid phases.

Also [1995Feu] calculated the thermodynamic functions of the ternary liquid using an association model

and compared with their measurements. A complete dataset for thermodynamic calculation of the whole

ternary system was reported by [2002Gor]. However, the reported dataset seems to contain errors more

215


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severe than a single typing error. An attempt to recalculate the published diagrams from this dataset resulted

in diagrams significantly different from the published ones.

A tentative set of ternary terms for the Gibbs-energies of liquid and the ternary solid solutions of the binary

intermediate phases is given in Table 3. It has to be used together with the three binary assessments from

the COST 507 action [1998Ans]. The fictitious term for the Ni2Mg type phase in the Cu-Mg system (Cu

occupation on Ni sites) is taken from the assessment of Cu-Mg-Zn [1998Ans]. All the interaction

parameters for the Laves phases tentatively are set independent on the occupation of the other sublattice,

thus, except the Cu-Ni interaction parameters, they were already evaluated in the binary assessments. The

ternary parameter for liquid stems from a transformation of Toop’s formula, modified by Hillert, into the

Muggianu formalism. Calculations by this dataset reproduce fairly well the all experimental points in the

Mg-rich part (> 50 at.% Mg) of the system, and may be taken as good approximations in the Mg-poor part.

A generalization of the Miedema model for the estimation of formation enthalpies of ternary and

higher-order intermetallics was developed by [1996Gon] and was successfully tested with respect to the

experimental data for alloys MgCu2-xNix. The estimated enthalpy values increased to some extent with

increasing Ni content.

A general solution model for the prediction of ternary thermodynamic properties from the binary

subsystems was proposed by [1995Cho] and tested successfully for several alloys of the Cu-Mg-Ni system.

Another such model, called parabolic model, was constructed by [1997Gan] and also tested successfully at

the Cu-Mg-Ni system.

References

[1939Vos] Vosskuehler, H., “Metallography of Magnesium and its Alloys” (in German), in

“Magnesium and its Alloys”, Beck, A., (Ed.), Springer Verlag, Berlin, 96 (1939) (Equi.

Diagram, Review, 1)

[1949Jae] Jaenecke, E., Short Reviewed Handbook of All Alloys, (in German), Carl Winter -

Universitaetsverlag, Heidelberg, 466-467 (1949) (Equi. Diagram, Review, 2)

[1951Koe] Koester, W., “Copper-Nickel-Magnesium Ternary System” (in German), Z. Metallkd.,

42(11), 326-327 (1951) (Equi. Diagram, Experimental, 4)

[1952Lie] Lieser, K.H., Witte, H., “Investigation of the Ternary Systems: Magnesium-Copper-Zinc,

Magnesium-Nickel-Zinc, and Magnesium-Copper-Nickel” (in German), Z. Metallkd.,

43(11), 396-401 (1952) (Equi. Diagram, Experimental, *, 17)

[1956Mik1] Mikheeva, V.I., Babayan, G.G., “The Melting Diagram of the Magnesium-Copper-Nickel

System” (in Russian), Dokl. Akad. Nauk SSSR, 108(6), 1086-1087 (1956) (Equi. Diagram,

Experimental, *, 6)

[1956Mik2] Mikheeva, V.I., Babayan, G.G., “About the Chemical Nature of the Ternary Intermetallic

Phases in the Systems Magnesium-Copper-Zinc and Magnesium-Copper-Nickel” (in

Russian), Dokl. Akad. Nauk SSSR, 109(4), 785-786 (1956) (Equi. Diagram,

Experimental, 8)

[1972Feh] Fehrenbach, P.J., Kerr, H.W., Niessen P., “The Constitution of Cu-Ni-Mg Alloys”,

J. Mater. Sci., 7(10), 1168-1174 (1972) (Equi. Diagram, Experimental, *, 8)

[1972Kom1] Komura, Y., Nakaue, A., “Crystal Structure of a New Stacking Variant of Friauf-Laves

Phases in the System Mg-Cu-Ni”, Acta Crystallogr., B28(3), 727-732 (1972) (Equi.

Diagram, Crys., Structure, Experimental, *, 13)

[1972Kom2] Komura, Y., Mitarai, M., Nakaue, A., Tsujimoto, S., “The Relation between Electron

Concentration and Stacking Variants in the Alloy Systems Mg-Cu-Ni, Mg-Cu-Zn and

Mg-Ni-Zn”, Acta Crystallogr., B28(3), 976-978 (1972) (Crys. Structure, Review, 12)

[1972Pre] Predel, B., Ruge, H., “About the Bond Relations in Laves Phases” (in German), Mater. Sci.

Eng., 9(6), 333-338 (1972) (Thermodyn., Experimental, 13)

[1974Kri] Kripyakevich, P.I., Melnik, P.I., “New Results on the Crystal Chemistry of Multilayer

Laves Phases” (in Russian), Akad. Nauk Ukr. SSR, Metallofizika, (52), 71-75 (1974) (Crys.


216


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Cu–Mg–Ni

[1977Ray] Raynor, G.V., “Constitution of Ternary and Some More Complex Alloys of Magnesium”,

Int. Met. Rev., 22, 65-96 (1977) (Equi. Diagram, Review, 83)

[1979Cha] Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Phase Diagrams and

Thermodynamic Properties of Ternary Copper-Metall Systems”, INCRA Monograph VI,


[1979Dri] Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L.,

Turkina, N.I., Binary and Multicomponent Copper-Base Systems (in Russian), Nauka

Moskow, 161-163 (1979) (Review, Equi. Diagram, 7)

[1983Dar] Darnaudery, J.P., Pezat, M., Darriet, B., “Influence of the Substitution of Nickel by Copper

in NiMg2 on the Hydrogen Storage” (in French), J. Less-Common Met., 92(2), 199-205

(1983) (Equi. Diagram, Experimental, *, 7)

[1983Kar] Karonik, V.V., Guseva, V.V., Ivanishev, A.V., Kolesnichenko, V.E., “Investigation of the

Mg-Ni-Cu and Mg-Ni-Ag Phase Diagram” (in Russian), Izv. Akad. Nauk SSSR, Met., (5),

220-226 (1983) (Equi. Diagram, Experimental, *, 7)

[1986She] Panwen, S., Yunshi, Z., Song, Z., Xianbao, F., Huatang, Y., Shengchang, C., “Chemical

Synthesis of Hydrogen-Storing Alloys (III) - Replacement-Diffusion Method for

Mg2Ni0.75Cu0.25”, Hydrogen Energy Progress VI, Proc. 6th World Hydrogen Energy Conf.,

Vienna, Austria, 2, 831-837 (1986) (Equi. Diagram, Crys. Structure, Experimental, 4)

[1991Gna] Gnansekaran, T., Ipser, H., “Thermodynamic Properties of Ternary Cu-Mg-Ni Alloys along

two Isopleths with x(Cu)/x(Ni) = 2.0 and 0.5”, COST 507 Leuven Proceedings; Part A,

Project A1, 1-16 (1991) (Experimental, Thermodyn., 25)

[1993Gna1] Gnanasekaran, T., Ipser H., “The Isopiestic Method Applied to an Investigation of the

Thermodynamic Properties of Ternary Cu-Ni-Mg Alloys”, J. Chim. Phys., 90(2), 367-372

(1993) (Experimental, Thermodyn., 16)

[1993Gna2] Gnanasekaran, T., Ipser, H., “Partial Thermodynamic Properties of Magnesium in Ternary

Liquid Copper-Magnesium-Nickel Alloys”, J. Non-Cryst. Solids, 156-158 (PT.1), 384-387

(1993) (Experimental, Thermodyn., Equi. Diagram, 12)

[1994Gna] Gnanasekaran, T., Ipser, H., “Thermodynamic Properties of Ternary Liquid Cu-Mg-Ni

Alloys”, Metall. Mater. Trans. B, 25(1), 63-72 (1994) (Thermodyn., Experimental, Equi.

Diagram, 25)

[1995Cho] Chou, K.-C., “A General Solution Model for Predicting Ternary Thermodynamic

Properties”, Calphad, 19(3), 315-325 (1995) (Thermodyn., 23)

[1995Feu] Feufel, H., Sommer, F., “Thermodynamic Investigations of Binary Liquid and Solid Cu-Mg

and Mg-Ni Alloys and Ternary Liquid Cu-Mg-Ni Alloys”, J. Alloys Compd., 224, 42-54

(1995) (Thermodyn., Experimental, Theory, 48)

[1995Ips] Ipser, H., Gnanasekaran, S., Boser, S., Mikler, H., “A Contribution to the Ternary

Copper-Magnesium-Nickel Phase Diagram”, J. Alloys Compd., 227, 186-192 (1995) (Equi.

Diagram, Experimental, *, 16)

[1995Jac] Jacobs, M.H.G., Spencer, P.J., “Thermodynamic Evaluation of the Systems Al-Si-Zn and

Cu-Mg-Ni”, J. Alloys Compd., 220, 15-18 (1995) (Thermodyn., 36)

[1996Gon] Goncalves, A.P., Almeida, M., “Extended Miedema Model: Predicting the Formation

Enthalpies of Intermetallic Phases with More than Two Elements”, Physica B, 228(3/4),

289-294 (1996) (Thermodyn., Theory, 19)

[1997Gan] Ganesan, R., Vana Varamban, S., “A Parabolic Model to Estimate Ternary Thermodynamic

Properties from the Corresponding Binary Data”, Calphad, 21(4), 509-519 (1997)

(Thermodyn., 19)

[1998Ans] Ansara, I., “Systems Cu-Mg, Cu-Ni, Mg-Ni” in “COST 507, Thermochemical Database for

Light Metal Alloys”, Ansara, I., Dinsdale A.T., Rand, M.H. (Eds.), European Communities,

Luxembourg, 1998, Vol.2, 170-174 (Cu-Mg), 175-177 (Cu-Ni), 218-220 (Mg-Ni), (Equi.

Diagram, Thermodyn., Assessment, 0)

217


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Cu–Mg–Ni

[1998Tsu] Tsushio, Y., Akiba, E., “Hydrogenation Properties of Mg-based Laves Phase Alloys”,

J. Alloys Comp., 269, 219-223 (1998) (Experimental, 19)

[2002Gor] Gorsse, S., Shiflet, G.J., “A Thermodynamic Assessment of the Cu-Mg-Ni Ternary

System”, Calphad, 26(1), 63-83 (2002) (Assessment, Calculation, Thermodyn., 35)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Ni,Cu)

Cu

< 1084.87

Ni

< 1455

cF4

Fm3m

Cu

a = 361.48

a = 352.40

Complete solid solution

pure Cu at 25°C [Mas2, V-C]

pure Ni at 25°C [Mas2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

pure Mg at 25°C [Mas2]

(NixCu1-x)2Mg

Cu2Mg

< 797

cF24

Fd3m

Cu2Mg

a = 692.8

a = 704.8

0 x 0.45 at 930°C [1952Lie]

at x = 0.4 [1952Lie]

at x = 0 [Mas2, V-C]

CuMg2

< 568

oF48

Fddd

CuMg2

a = 907.0

b = 528.4

c = 1825.0

[Mas2, V-C]

(Ni1-xCux)2Mg

Ni2Mg

< 1147

hP24

P63/mmc

Ni2Mg

a = 486.1

c = 1594

a = 482.4

c = 1582.6

0 x 0.49 at 930°C [1952Lie]

at x = 0.39 [1952Lie, V-C]


(Ni1-xCux)Mg2

NiMg2

< 760

hP18

P6222

NiMg2

a = 525

c = 1355

a = 519.8

c = 1321

0 x 0.85 at 600°C [1983Dar]

at x = 0.85 [1983Dar]


* (Ni1-xCux)2Mg

at least < 800

hP36

P63/mmc

(Ni1-xCux)2Mg

a = 491.7

c = 2404.0

0.5 < x < 0.55

[1972Kom1, V-C]

218


MSIT®

Cu–Mg–Ni


* Values given in parentheses are uncertain by several at.%.

Table 3: Ternary Parameters for the Cu-Mg-Ni System. To be Used Together with the Binary Parameter

Datasets Cu-Mg, Cu-Ni and Mg-Ni of the COST 507 Action [1998Ans]

Reaction T [°C] Type Phase Composition* (at.%)

Cu Mg Ni

L + Ni2Mg Cu2Mg 930 p2

max

L

Ni2Mg

Cu2Mg

50.0

32.7

36.7

33.3

33.3

33.3

16.7

34.0

30.0

L + Ni2Mg Cu2Mg + (Ni,Cu) 808 U1 L

Ni2Mg

Cu2Mg

(Ni,Cu)

(71)

29

42

(78)

19

31

31

(3)

(10)

40

27

(19)

L + Ni2Mg Cu2Mg + NiMg2 658 U2 L

Ni2Mg

Cu2Mg

NiMg2

33

24

42

(11)

58

34

34

67

9

42

24

(22)

L NiMg2 + CuMg2 567 e3

max

L

NiMg2

CuMg2

32.3

26.5

32.5

66.7

66.7

66.7

1.0

6.8

0.8

L + NiMg2 Cu2Mg + CuMg2 553 U3 L

Cu2Mg

NiMg2

CuMg2

39

63

26

32

60

35

67

67

1

2

7

1

L (Mg) + NiMg2 + CuMg2 480 E1 L

(Mg)

NiMg2

CuMg2

14

0.013

25

32

84

100

67

67

2

0

8

1

Parameter T-range [K] Value

LCu,Mg,Niliq 298-6000 +7500. -9.2 T

0GMg:NiLaves-C15 - 0GMg:Ni

Laves-C36 298-6000 +4000

0GNi:MgLaves-C15 - 0GNi:Mg

Laves-C36 298-6000 -4000

0L*:Cu,MgLaves-C15 298-6000 +13011.

0LCu,Mg:*Laves-C15 298-6000 +6599.

0L*:Cu,NiLaves-C15 298-6000 +25100. -8.0 T

0LCu,Ni:*Laves-C15 298-6000 +25100. -8.0 T

0L*:Mg,NiLaves-C15 298-6000 +50000.

0LMg,Ni:*Laves-C15 298-6000 +50000.

219


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Cu–Mg–Ni

0GMg:CuLaves-C36 - 0GMg:Cu

Laves-C15 298-6000 +4000.

0GCu:MgLaves-C36 - 0GCu:Mg

Laves-C15 298-6000 -4000.

0L*:Cu,MgLaves-C36 298-6000 +13011.

0LCu,Mg:*Laves-C36 298-6000 +6599.

0L*:Cu,NiLaves-C36 298-6000 +25100. -8.0 T

0LCu,Ni:*Laves-C36 298-6000 +25100. -8.0 T

0L*:Mg,NiLaves-C36 298-6000 +50000.

0LMg,Ni:*Laves-C36 298-6000 +50000.

0GNi:MgLaves-C36 - 0GNi

SER - 2 0GMgSER 298-6000 -30000. +8.0 T

0GCu:MgNiMg2 - 0GCu

SER - 2 0GMgSER 298-6000 -26000. +0.5 T

Parameter T-range [K] Value

10 20 30 40 50 60

700

800

900

1000

1100

1200

Cu 0.00

Mg 33.30

Ni 66.70

Cu 66.70

Mg 33.30

Ni 0.00Cu, at.%

Te

mp

era

ture

, °C

Ni2Mg

1147

L

930 p2

797

Cu2Mg(Ni0.45Cu0.55)2Mg

Fig. 1: Cu-Mg-Ni.

The pseudobinary

system

Ni2Mg-Cu2Mg

220


MSIT®

Cu–Mg–Ni

Fig. 2: Cu-Mg-Ni. Reaction scheme

l (Cu)+Cu2Mg

725 e2

l Cu2Mg+CuMg

2

552 e3

l (Mg)+CuMg2

485 e6

l (Ni)+Ni2Mg

1097 e1

l+Ni2Mg NiMg

2

760 p2

l (Mg)+NiMg2

506 e5

L+Ni2Mg Cu

2Mg

930 p1max

L+Ni2Mg Cu

2Mg+(Ni,Cu)808 U

1

L+Ni2Mg Cu

2Mg+NiMg

2658 U

2

L+NiMg2

Cu2Mg+CuMg

2553 U

3

L (Mg)+NiMg2+CuMg

2480 E

1

L+Cu2Mg+NiMg

2

Ni2Mg+Cu

2Mg+(Ni,Cu)

Ni2Mg+Cu

2Mg+NiMg

2

Cu2Mg+NiMg

2+CuMg

2

(Mg)+NiMg2+CuMg

2

L+NiMg2+CuMg

2

567 e3

Cu-Mg Mg-NiCu-Mg-Ni

20

40

60

80

20 40 60 80

20

40

60

80

Ni Cu


Axes: at.%

e5

E1

e6

e3

CuMg2

U3

e4

Cu2Mg

U1

e2

(Ni,Cu)

e1

Ni2Mg

NiMg2

U2

p1

p2

Fig. 3: Cu-Mg-Ni.

Calculated projection

of the four-phase

equilibrium planes

and lines of double

saturation of liquidus

and solidus

221


MSIT®

Cu–Mg–Ni

20

40

60

80

20 40 60 80

20

40

60

80

Ni Cu


Axes: at.%

14001350 1300

1250 1200

1150 1100

1100 1050 1000 950 900

Ni2Mg

e1

850 800

U1

600°C550

550 600

650700

NiMg2

p2

(Mg)

E1

CuMg2

e3

U3

e4

Cu2Mg

e2

(Ni,Cu)

p1

U2

e5

1000

950900

850800

1050

Fig. 4: Cu-Mg-Ni.

Calculated liquidus

surface

20

40

60

80

20 40 60 80

20

40

60

80

Ni Cu


Axes: at.%

NiMg2 CuMg

2

Cu2Mg

(Ni,Cu)

(Ni,Cu)+Ni2Mg

(Ni,Cu)+Ni2Mg+Cu

2Mg

Ni2Mg+Cu

2Mg+NiMg

2

Ni2Mg+NiMg

2

Ni2Mg

Cu2Mg+CuMg

2

(Ni,Cu)+

Cu2Mg

Fig. 5: Cu-Mg-Ni.


section at 475°C

222


MSIT®

Cu–Mg–Ni

70 60

700

800

Cu 76.90

Mg 23.10

Ni 0.00

Cu 56.20

Mg 18.80

Ni 25.00Cu, at.%

Te

mp

era

ture

, °C

e2

U1 808

(Ni,Cu)+Ni2Mg+Cu2Mg

L+(Ni,Cu)+Cu2Mg

L+Ni2Mg+Cu2Mg

(Ni,Cu)+Cu2Mg

Fig. 6: Cu-Mg-Ni.

Isopleth along the line

of double saturation

of liquid with respect

to (Ni,Cu) and Cu2Mg

10 20 30 40

500

750

1000

1250

Mg Cu 50.00

Mg 0.00

Ni 50.00Cu, at.%

Te

mp

era

ture

, °C

L

L+(Mg)

L+Ni2Mg+NiMg2

NiMg2

L+Cu2Mg+NiMg2

L+NiMg2

L+Ni2Mg

(Ni,Cu)+Ni2Mg

L+(Ni,Cu)

L+(Ni,Cu)+Ni2Mg

658

808

(Mg)+NiMg2

Ni2Mg+Cu2Mg+

NiMg2

Cu2Mg+Ni2Mg

930

Fig. 7: Cu-Mg-Ni.

Calculated isopleth at

Cu:Ni=1:1

223


MSIT®

Cu–Mg–Ni

10 20

400

500

600

700

800

Cu 29.00

Mg 71.00

Ni 0.00

Cu 0.00

Mg 71.00

Ni 29.00Ni, at.%

Te

mp

era

ture

, °C

L

L+NiMg2+CuMg2

(Mg)+NiMg2

L+(Mg)+NiMg2

L+CuMg2 L+NiMg2

480

(Mg)+CuMg2

(Mg)+NiMg2+CuMg2

L+(Mg)+CuMg2

Fig. 8: Cu-Mg-Ni.

Isopleth at 71 at.%

Mg

224


MSIT®

Cu–Mg–Si

Copper – Magnesium – Silicon

Nataliya Bochvar, Evgeniya Lysova and Lazar Rokhlin

Literature Data

Many research groups dealt with the constitution of the Cu-Mg-Si system [1933Por, 1936Lav, 1938Wit,

1939Wit, 1942Lav, 1954Kle, 1953Ber, 1953Nag, 1956Ber, 1960Asc, 1984Kom, 1984Mat, 1985Far,

1986Mat, 1987Mat]. Reviews on phase equilibria were given by [1939Vos, 1949Jae, 1952Pie, 1977Dri].

Reviews on crystal structures were given by [1959Ray, 1968Kry, 1969Tes]. Agreement exists on the

formation of three ternary compounds; Cu1.5MgSi0.5 (Cu3Mg2Si), Cu16Mg6Si7 and

(Cu0.8Si0.2)2(Mg0.88Cu0.12), which essentially determine phase equilibria.

[1933Por] studied the region Mg-Mg2Si-Cu1.5MgSi0.5-Cu2Mg by thermal analysis and optical microscopy.

Based on the melting temperature of 927°C for the ternary compound Cu1.5MgSi0.5, [1933Por] claimed the

existence of two “pseudobinary sections”, Mg2Si-Cu1.5MgSi0.5 and CuMg2-Cu1.5MgSi0.5 with eutectic

points at 857 and 565°C, respectively. Two invariant four-phase equilibria at 508°C (transition type) and at

479°C (eutectic type) were reported, as well.

[1936Lav, 1942Lav] revealed connection between concentration of valence electrons and crystal structure

type formed along the isopleth section at 33.3 at.% Mg.

[1938Wit, 1939Wit] investigated the vertical section at 33.3 at.% Mg up to 40 at.% Si by X-ray diffraction

and thermal analysis. Peritectic formation of the compound Cu1.5MgSi0.5 and its homogeneity range were

established and furthermore a polymorphic transformation between 870 and 890°C: the high temperature

form adopts the Ni2Mg type, whilst the low temperature form is Cu1.5MgSi0.5 type as an ordered version of

the MgZn2 type. The lattice parameters of the compound Cu1.5MgSi0.5 were given by [1938Wit, 1939Wit,

1970Sch]. The lattice parameters of the Cu2Mg-base solid solution were measured by [1938Wit, 1939Wit,

1960Asc, 1979Ell]. A second ternary compound was discovered at Cu16Mg6Si7 [1938Wit, 1939Wit]. The

details of its crystal structure were studied by [1953Ber, 1953Nag, 1956Ber].

[1984Kom, 1984Mat, 1986Mat, 1987Mat] found the third ternary compound near the composition

Mg(Cu0.8Si0.2)2.4 with a new ordered by low symmetry variant of the Cu2Mg type. These investigations

were carried out using the alloys of different compositions within 25-35 at.% Mg, 10-20 at.% Si, rest Cu.

The alloys were annealed at 500°C for 10 days and small single crystals suitable for X-ray diffraction

studies were obtained by crushing. [1986Mat, 1987Mat] determined also the homogeneity range of this new

ternary phase. In the same investigations [1986Mat, 1987Mat] the homogeneity range of the binary

compound Cu2Mg was established to extend into the ternary system up to 25-33.3 at.% Mg at about 13 at.%

Si. These data correspond to the results of [1954Kle] who determined from susceptibility measurements the

boundary of the Cu2Mg homogeneity along the 33.3 at.% Mg section to be more precisely at 13.3 at.% Si.

Phase equilibria in the Cu corner of the Cu-Mg-Si system were investigated in detail using thermal analysis,

optical microscopy, X-ray diffraction and chemical analysis on 250 alloys [1960Asc]. As a result of the

study liquidus surface and isothermal section at 450°C were constructed for the Cu corner of the phase

diagram. [1960Asc] established a series of three-phase and four-phase invariant equilibria in the studied part

of the system. According to [1960Asc], the ternary compound Cu16Mg6Si7 formed from melt by a

four-phase peritectic reaction at 826°C; the melting point of the ternary compound Cu1.5MgSi0.5 was

determined to be 930°C as compared with 927°C [1933Por]. [1960Asc] also proposed the liquidus surface

for the rest of the ternary system based on the experimental data of [1933Por, 1938Wit, 1939Wit] and some

own experimental results.

[1985Far] studied the alloy containing 34.7Cu-27.5Mg-37.8Si (at.%) by DTA, X-ray diffraction, optical

microscopy and SEM. According to [1985Far], the alloy represents the ternary eutectic

Cu16Mg6Si7+Mg2Si+(Si) with melting point of 770°C. These results, however, are in contradiction to

[1960Asc] who showed for the same composition the ternary eutectic Cu1.5MgSi0.5+Mg2Si+(Si) and to

[1938Wit, 1939Wit], who showed the eutectic temperature to be below 765°C.

225


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Cu–Mg–Si

Binary Systems

The binary system Cu-Mg is accepted from [2002Iva]. The binary system Mg-Si is accepted from

[1984Nay]. The binary system Cu-Si is accepted from [2002Leb].

Solid Phases

Three ternary compounds are found in the system Cu-Mg-Si. The ternary compound 1, Cu1.5MgSi0.5 is

formed by a three-phase peritectic reaction at 930°C. It exists in two crystal forms with a transition

temperature of about 870 to 890°C and exhibits a homogeneity range 16.7-23.3 at.% Si along the isopleth

for 33.3 at.% Mg [1938Wit, 1939Wit, 1960Asc]. The ternary compound 2, Cu16Mg6Si7 is formed by the

four-phase peritectic reaction at 826°C with a limited homogeneity range. Deviation from the stoichiometric

composition amounts to about 0.5 at.% [1960Asc]. The ternary compound 3 with approximate composition

Mg(Cu0.8Si0.2)2.4 exists in solid state at 500°C. The homogeneity range of this ternary compound is within

the limits 25.5 to 30.0 at.% Mg, 16 to13.5 at.% Si, rest Cu [1986Mat, 1987Mat]. The highest and lowest

temperature of existence and reaction of its formation are unknown. The binary phase Cu2Mg dissolves up

to 13.3 at.% Si along the isopleth for 33.3 at.% Mg [1954Kle]. The homogeneity range of this phase at

500°C enlarges from 32.5-35 at.% Mg in the binary system Cu-Mg to 25-33.3 at.% Mg at 13 at.% Si in the

ternary system [1986Mat, 1987Mat]. Dissolution of Si in Cu2Mg significantly increases its melting

temperature up to 950°C at 13.3 at.% Si [1938Wit, 1939Wit]. The binary copper silicides '', 1 and

dissolve at 450°C up to 1.3, 5.3 and 4.3 at.% Si, respectively [1960Asc]. Crystal structure and the lattice

parameters of all solid phases pertinent to the ternary system are presented in Table 1.


There is no pseudobinary section in the Cu-Mg-Si phase diagram, although the sections 1-Mg2Si,

1-CuMg2 were considered to be of pseudobinary nature by [1933Por] and the sections 1-(Si), 2-(Si) and

2- were considered pseudobinary by [1960Asc]. These viewpoints are wrong, however, as both ternary

compounds 1 and 2 melt incongruently.


In the system Cu-Mg-Si 10 three-phase and 18 four-phase invariant equilibria were established with

participation of liquid phase. The partial reaction scheme is shown in Figs. 1a, 1b. Compositions of the

phases participating in the invariant equilibria are presented in Table 2. Three-phase invariant equilibria are

as follows, L 1+Mg2Si (e2max), L + (e3max), L 1+(Si) (e7max), L 2+ (e8max), L (Cu)+ (e9max),

L 2+(Si) (e10max), L + (e12max), L 1+ (e13max), L 1+CuMg2 (e16max), L+Cu2Mg 1 (p1max).

Temperatures of the eutectic points e7max and e10max are not determined experimentally, but they can be

estimated. Accordingly, the e7max temperature is to be higher than 765°C. This temperature is shown in the

vertical section for the constant 33.3 at.% Mg [1938Wit, 1939Wit] for the point corresponding to its

intersection with the monovariant eutectic line L 1+(Si).

The e10max temperature should be higher than 739°C in correspondence to the temperature E2.

Temperatures of all four-phase invariant equilibria, except E1 and U8, are adopted from the experiments

[1960Asc]. The temperature of the E1 reaction is estimated to be below 765°C from comparison with the

vertical section at 33.3 at.% Mg. Considering the reaction scheme U8 must be between 565 (e16max) and

552°C (e17).

The only invariant equilibrium in solid state established by experiments [1960Asc] is the eutectoid

decomposition of at 609°C. Nevertheless, this reaction can not be recognized as a eutectoid one because

of contradiction with the decomposition reaction of in the binary Cu-Si system. Therefore, the reaction at

609°C connected with the decomposition of is assumed to be of the transition type U7 as shown in the

reaction scheme (Fig. 1). Taking into consideration the established invariant equilibria in the ternary system

and the accepted binary system Cu-Si, it is reasonable to propose the existence of three invariant solid state

equilibria in solid state U3, p5min and p7min. The equilibria are shown in Fig. 1.

226


MSIT®

Cu–Mg–Si

Liquidus Surface

The projection of the liquidus surface is shown in Fig. 2. It is redrawn from [1960Asc] replacing mass% for

at.%. Fig. 3 presents the liquidus surface of the copper corner in enlarged view. A remarkable feature of the

liquidus surface is the existence of the primary crystallization fields of the phases , and which are

formed in the Cu-Si binary by peritectoid reactions. The field of the primary crystallization of the ternary

compound 3 is absent and there is no reason for its formation from the liquid. The liquidus isotherms in the

Cu corner of the system are drawn after [1960Asc]. Other isotherms are assumed taking only into account

the corresponding binary systems.

Isothermal Sections

Figure 4 shows fragments of the isothermal section of the phase diagram at 500°C. The existence areas of

the solid phases 3, and adjoining them two-phase and three-phase areas are drawn using experimental

data presented in [1986Mat, 1987Mat]. The region of existence for the solid phases 1 and 2 are drawn

according to [1938Wit, 1939Wit]. Figure 5 shows isothermal section of the phase diagram at 450°C. It is

drawn for the Cu corner after [1960Asc]. Moreover, the field is corrected according to the data [1986Mat,

1987Mat] including also 3. The rest of the section is tentatively constructed from the binary phase diagrams

[2002Iva, 2002Leb] and results of [1933Por].


The partial vertical section at 33.3 at.% Mg is shown in Fig. 6. It is based on the data [1938Wit, 1939Wit,

1954Kle] with a slight correction for the existence of the three-phase area L+ 1+(Si).

Thermodynamics

[1996Gan] determined thermodynamic activities and partial molar enthalpies of magnesium for liquid

phase in the temperature range 740-1050°C by measuring magnesium vapor along isopleth with a copper

to silicon ratio 7:3.

The heat of fusion (eutectic transformation) was measured by [1985Far] for the ternary eutectic at

Cu34.67Mg27.515Si37.82 to be Hfus = 16.61 kJ mol-1 employing DSC.

[1997Ips] determined enthalpy of mixing along an isopleth with a constant concentration ratio of xCu/xSi =

7/3 by solution calorimetry method. Results of the experiments are shown in Fig. 7. Moreover, [1997Ips]

determined magnesium activity as function of composition along the same isopleth xCu/xSi = 7/3 at 900°C

(Fig. 8). Magnesium activity data were derived from the magnesium vapor pressure measurements.

Miscellaneous

Solubility of hydrogen in the Laves phases along the section at 33.3 at.% Mg and various thermodynamic

characteristics (enthalpy and entropy) were determined by [1957Wit, 1957Lie, 1957Sie]. The solubility of

hydrogen at 0.1 MPa generally increases in the alloys from 450 to 550°C [1957Lie, 1957Sie]: hydrogen

isotherms for copper-rich alloys reveal a nonlinear behavior with a minimum around the composition

64.7Cu-33.3Mg-2Si (at.%) and a flat maximum around the composition 58.7Cu-33.3Mg-8Si (at.%)

[1957Lie, 1957Sie]. [1957Lie] determined the entropy change for a transition of H2 gas dissolved in the

alloys.

[1993Mur] used the ternary Cu-Mg-Si phase diagram for analysis and description of the Al-rich part of the

quaternary Al-Cu-Mg-Si system.

References

[1933Por] Portevin, A., Bonnot, M., “Contribution to Study of the Constitution of the Ternary

Magnesium-Copper-Silicon Alloys” (in French), Compt. Rend. Acad. Sci. Paris, 196,


227


MSIT®

Cu–Mg–Si

[1936Lav] Laves, F., Witte, H., “Influence of Valence Electron to Crystal Structure of Ternary

Magnesium Alloys” (in German), Metallwietschaft, 15, 840-842 (1936)

(Crys. Structure, 10)

[1938Wit] Witte, H., “The Study of the Crystal Chemistry of Alloys. II. Investigations in the System

Magnesium-Copper-Silicon with Special Reference to the Section MgCu2-MgSi2” (in

German), Z. Angew. Mineral., 1, 255-268 (1938) (Equi. Diagram, Experimental., #, 12)

[1939Vos] Vosskuehler, H., “Metallography of Magnesium and its Alloys” (in German), in

“Magnesium und Seine Legierungen”, Beck, A. (Ed.), Berlin, 96 (1939) (Equi. Diagram,

Review, 1)

[1939Wit] Witte, H., “The Study of the Crystal Chemistry of Alloys: Investigations in the System

Magnesium-Copper-Silicon with Special Reference to the Section MgCu2-MgSi2” (in

German), Metallwirtschaft, 18(22), 459-463 (1939) (Equi. Diagram, Experimental, *, #, 12)

[1942Lav] Laves, F., Wallbaum, H.J., “On the Influence of Geometrical Factors on the

Stoichiometrical Formula of Metallic Bonds Demonstrated of Crystal Structure of KNa2”,

(in German), Z. Anorg. Allg. Chem., 250, 110-120 (1942) (Crys. Structure, Experimental, 9)

[1949Jae] Jaenecke, E., “Cu-Mg-Si” (in German), Kurzgefasstes Handbuch aller Legierungen,

577-578 (1949) (Equi. Diagram, Review, 2)

[1952Pie] Pietsch, E.H.E., Meyer, R.J., “Magnesium-Copper-Silicon” (in German), Gmelins

Handbuch der Anorg. Chemie, Verlag. Chemie, GmbH., Weiheim/Bergstasse, 27(A4),

714-716 (1952) (Equi. Diagram, Reviev, *, 5)

[1953Ber] Bergman, G., Waugh, J.L.T., “The Crystal Structure of the Intermetallic Compound

Mg6Si7Cu16”, Acta Crystallogr., 6(1), 93-94 (1953) (Crys. Structure, Experimental, 3)

[1953Nag] Nagorsen, G., Witte, H., “The Crystal Structure of Mg6Si7Cu16” (in German), Z. Anorg.

Allg. Chem., 271, 144-149 (1953) (Crys. Structure, Experimental, 3)

[1954Kle] Klee, H., Witte, H., “The Magnetic Susceptibility of Ternary Magnesium Alloys and its

Explanation with Point of View of Electronic Theory of Metals”, Z. Phys. Chem. (Leipzig),

202 , 352-377 (1954) (Equi. Diagram, Experimental, #, 30)

[1956Ber] Bergman, G. ans Waugh, J.L.T., “The Crystal Structure of the Intermetallic Compound

Mg6Si7Cu16”, Acta Crystallogr., 9(3), 214-217 (1956) (Crys. Strycrure, Experimental, 10)

[1957Lie] Lieser, K.H., Witte, H., “Solubility of Hydrogen in Alloys. IV. Discussion” (in German),

Z. Elektrochem., 61(3), 367-376 (1957) (Experimental, 31)

[1957Sie] Siegelin, W., Lieser, K.H., Witte, H., “Solubility of Hydrogen in Alloys. III. Study of

Ternary MgCu2-MgAl2, MgCu2-MgSi2, MgNi2-MgCu2 Systems and Binary Ag-Cd

System” (in German), Z. Elektrochem., 61(3), 359-366 (1957) (Experimental, 9)

[1957Wit] Witte, H., “Solubility of Hydrogen in Alloys” (in German), Neue Huette, 2(12), 749-756


[1959Ray] Raynor, G.V., The Physical Metallurgy of Magnesium amd Its Alloys, London, New York,

Paris, Los Angeles: Pergamon Press, 531 p. (1959) (Crys. Structure, Review, 35)

[1960Asc] Aschan, L. J., “Studies on the Ternary System Cu-Mg-Si”, Acta Polytech. Scand., 11(285),

1-63 (1960) (Equi. Diagram, Experimental, *, #, 43)

[1968Kry] Krypyakevich, P.I., Gladyshevskii, E.I., Cherkashin, E.E., “Problems of Crystalchemistry

of Intermetallic Compounds in Papers of Co-workers of the Department of Inorganic

Chemistry of the L,viv University” (in Ukrainian), Visnik L'viv Univ., Ser. Khim, (10), 90-99

(1968) (Crys. Structure, Reviev, 115)

[1969Tes] Teslyuk, M.,Yu., Intermetallic Compounds with Structure of Laves Phases (in Russian),

Moskow, Nauka, 136p. (1969) (Crys. Structure, Review, Theory, 312)

[1970Sch] Schuster, H.U., Bockelmann, W., Captuller, J., “Ternary Phases in the

Magnesium-Copper-Germanium System” (in German), Z. Naturforsch. B, 25B (11),

666-668 (1970) (Crys. Structure, 2)

[1977Dri] Drits, M.E., Bochvar, N.R., Kadaner, E.S., Padezhnova, E.M., Rokhlin, L.L., Sviderskaya,

Z.A., Turkina, N.I., “Magnesium-Silicon-Copper” (in Russian), in Phase Diagrams of

228


MSIT®

Cu–Mg–Si

Alunimium and Magnesium Alloys, Moscow, Nauka, 167-168 (1977) (Equi. Diagram,

Review, 6)

[1979Ell] Ellner, M., Predel, B., “Neutron Diffraction Investigation of Ternary Laves Phases with

MgCu2-Type” (in German), J. Solid State Chem., 30, 209-221 (1979) (Crys. Structure,

Experimental, 26)

[1984Kom] Komura, Y., Matsunaga, T., “A New Ordered Structure of the Off-Stoichiometric Laves

Phase Having C15-Type Structure in Mg-Cu-Si System”, Mater. Res. Soc. Symp. Proc., 21,


[1984Mat] Matsunaga, T., Koders, E., Komura, Y., “A New Ordered Structure of C15-Type Laves

Phase, Mg28.4Cu57.9Si13.7”, Acta Crystallogr., Sect. C: Cryst. Struct. Commun., 40C (10),


[1984Nay] Nayeb-Hashemi, A.A., Clark, J.B., “The Mg-Si (Magnesium-Silicon) System”, Bull. Alloy

Phase Diagrams, 5(6), 584-592 (1984) (Equi. Diagram, Review, 49)

[1985Far] Farkas, D., Birchenall, C.E., “New Eutectic Alloys and Their Heats of Transformation”,

Metall. Trans. A, 16A(3), 323-328 (1985) (Experimental, 18)

[1986Mat] Matsunaga, T., Komura, Y., “A New Ordered Phase of MgCu2-Type Structure in the

Mg-Cu-Si System” (in Japanese), Nippon Kinzoku Gakkai-Shi (J. Jpn. Inst. Met.), 50(7),

611-615 (1986) (Crys. Structure, Experimental, *, #, 19)

[1987Mat] Matsunaga, T., “A Study of New Ordered Structure in the Magnesium-Copper-Silicon

Ternary System”, J. Sci. Hiroshima Univ., Ser. A: Phys. Chem., 51(3), 247-275 (1987)

(Crys. Structure, Experimental, *, #, 28)

[1993Mur] Murray, J.L., “Industrial Applications of Multicomponent Aluminum Phase Diagrams”,

J. Chim. Phys., 90, 151-166 (1993) (Equi. Diagram, Experimental, 8)

[1996Gan] Ganesan, V., Ipser, H., “Partial Thermodynamic Properties of Magnesium in Ternary

Cu-Mg-Si Alloys”, J. Non-Cryst. Solids, 205-207(2), 711-715 (1996) (Experimental, 11)

[1997Ips] Ipser, H., Sommer, F., “Thermochemistry of Magnesium Based Light Alloys”, Proc.

-Electrochem. Soc., (High Temperature Materials Chemistry), 97-39, 31-37 (1997)

(Thermodyn., Experimental, 18)

[2002Iva] Ivanchenko, V., Ansara, I., “Cu-Mg (Copper-Magnesium)”, MSIT Binary Evaluation

Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International

Services GmbH, Stuttgart; Document ID: 20.10551.1.20, (2002) (Crys. Structure, Equi.

Diagram, Assessment, 13)

[2002Leb] Lebrun, N., Dobatkina, T., Kuznetsov, V., Li, C., “Cu-Si (Copper-Silicon)”, MSIT Binary





Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

pure Mg at 25°C

[2002Iva]

(Cu)

< 1084.62

cF4

Fm3m

Cu

a = 361.46

0 to 11.1 at.% Si

pure Cu at 25°C

[2002Leb]

229


MSIT®

Cu–Mg–Si

(Si)

< 1414

cF8

Fd3m

C (diamond)

a = 543.06 pure Si at 25°C

[2002Leb]

, Cu2-x-yMg1+ySix< 950

Cu2Mg

< 797

cF24

Fd3m

Cu2Mg a = 699.9

a = 701.1

a = 696.9

a = 698.3

a = 702.1

a = 701.3

0 x 0.39,

-0.25 y 0.05 [1938Wit, 1939Wit]

at x = 0.2, y = 0 [1938Wit, 1939Wit]

at x = 0.2, y = 0 [1979Ell]

at x = 0.23, y = 0 [1960Asc]

at x = 0.38, y = 0 [1938Wit, 1939Wit]

64.7 to 69 at.% Cu [2002Iva]

at x = 0, y = 0 [2002Iva]

at x = 0, y = 0 [1938Wit, 1939Wit]

CuMg2

< 568

oF48

Fddd

CuMg2

a = 904.4 0.1

b = 527.5 0.1

c = 1832.8 0.2

a = 907

b = 528.4

c = 1825

a = 905

b = 528.3

c = 1824.7

[2002Iva]

Mg2Si

< 1085

cF12

Fm3m

CaF2

a = 633.8 [1984Nay]

, Cu7Si

842 - 552

hP2

P63/mmc

Mg

a = 256.05

c = 418.46

11.05 to 14.5 at.% Si

at 12.75 at.% Si [2002Leb]

, ~Cu6Si

853 - 787

cI2

Im3m

W

a = 285.4

14.2 to 16.2 at.% Si [2002Leb]

at 14.9 at.% Si [2002Leb]

, Cu5Si (h)

824 - 711

t**

a = 881.5

c = 790.3

17.6 to19.6 at.% Si

sample was annealed at 700°C

[2002Leb]

1, Cu5Si (r)

< 729

cP20

P4132

Mn

a = 619.8 17.15 to 17.6 at.% Si [2002Leb]

, Cu15Si4< 800

cI76

I43d

Cu15Si4

a = 961.5 21.2 at.% Si [2002Leb]

, Cu3Si (h2)

859 - 558

hR*

R3m

or

t**

a = 247

= 109.74°

a = 726.7

c = 789.2

23.4 to 24.9 at.% Si [2002Leb]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

230


MSIT®

Cu–Mg–Si


', Cu3Si (h1)

620 - 467

hR*

R3m a = 472

= 95.72°

23.2 to 25.2 at.% Si [2002Leb]

'', Cu3Si(r)

< 570

o**

a = 7676

b = 700

c = 2194

23.3 to 24.9 at.% Si [2002Leb]

* 1, Cu1.5MgSi0.5 (h)

930 - 870

hP24

P63/mmc

MgNi2

[1938Wit, 1939Wit]

* 1, Cu1.5MgSi0.5 (r)

< 870

hP12

P63/mmc

Cu1.5MgSi0.5

(ordered

MgZn2)

a = 500.4 0.5

c = 787.3 0.8

a = 501.4

c = 788.8

a = 499.8

c = 795.1

at 16.7Si-33.3Mg-50.0Cu (at.%)

[1938Wit, 1939Wit]

at 16.7Si-33.3Mg-50.0Cu (at.%)

[1970Sch]

at 23.3Si-33.3Mg-43.4Cu (at.%)

[1938Wit, 1939Wit]

* 2, Cu16Mg6Si7< 826

cF116

Fm3m

Mn23Th6

a = 1167 1

a = 1165 2

[1938Wit, 1939Wit, 1960Asc]

exp = 5.66 Mg m-3

[1953Nag]

* 3, (Cu0.8Si0.2)2

(Mg0.88Cu0.12)

cP24

P4132

or

P4332

(ordered

derivative of

Cu2Mg)

a = 697.76 0.06 at 500°C,

exp = 5.66 Mg m-3 [1987Mat]


Cu Mg Si

L + 1 930 p1max L

1

48.7

53.4

50.0

33.3

33.3

33.3

18.0

13.3

16.7

L 1 + Mg2Si 857 e2max L

1

Mg2Si

36.7

~46.2

0

42.2

33.3

66.7

21.1

~20.5

33.3

L + 835 e3max L ~82.6 ~1.7 ~15.7

L + + 1 2 826 P1 L

1

2

54.9

53.4

~45.2

55.0

21.3

33.3

33.3

21.4

23.8

13.3

~21.5

23.6

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

231


MSIT®

Cu–Mg–Si

L + (Cu) + 824 U1 L ~83.0 ~2.1 ~14.9

L + + 1 806 P2 L 80.5 1.6 17.9

L + 1 + 800 U2 L 80.9 2.1 17.0

L + + 1 772 P3 L ~77.8 ~3.8 ~18.4

L + + 763 U4 L ~77.3 ~4.2 ~18.5

L + 2 ~770 e8max L ~65.9 ~10.5 ~23.6

L (Cu) + ~760 e9max L

(Cu)

~78.2

~91.6

53.4

~15.5

~5.2

33.3

~6.3

~3.2

13.3

L + 1 + 745 U5 L 76.2 5.6 18.2

L + 2 + (Si) 739 E2 L ~65.2 ~10.0 ~24.8

L + ~723.5 e12max L 66.4 14.7 18.9

L + + 2 723 E4 L 65.8 14.8 19.4

L + (Cu) + 718 U6 L 76.4 9.7 13.9

L 1 + 711 e13max L 75.5 8.5 16.1

L + 1 + 701 E5 L ~74.4 ~8.6 ~17.0

L + 1 + 693 E6 L 76.4 8.8 14.8

+ 1 + (Cu) 609 U7 ~82.8 ~3.5 ~13.7

L CuMg2 + 1 565 e16max L

CuMg2

1

33.1

33.3

~43.7

66.4

66.7

33.3

0.5

0

~23.0

L + 1 CuMg2 + Mg2Si 508 U9 L

1

CuMg2

Mg2Si

19.0

~43.7

33.3

0

80.55

33.3

66.7

66.7

0.45

~23.0

0

33.3

L (Mg) + CuMg2 + Mg2Si 479 E7 L

(Mg)

CuMg2

Mg2Si

15.6

~0

33.3

0

84.0

~100

66.7

66.7

0.4

~0

0

33.3


Cu Mg Si

232


MSIT®

Cu–Mg–Si

Fig

. 1a:

C

u-M

g-Si

. Par

tial

rea

ctio

n s

chem

e, p

art

1

Cu

-Si

Mg

-Si

Cu

-Mg-S

i

l M

g2S

i +

(S

i)

945.6

e 1

L +

γτ 1

93

0p

1m

ax

Lτ 1

+ M

g2S

i

85

7e 2

max

L+

γ +

τ 1τ 2

82

6P

1

l +

(C

u)

β8

52

p2

β +

(C

u)

χ8

42

p3

Lβ

+ χ

~8

35

e 3m

ax

L+

β (

Cu

) + χ

82

4U

1l +

βδ

82

4p

4

l η

+ δ

82

0e 4

(Cu

) + β

χ?

p5m

in

L +

β +

δ γ 1

80

6P

2

l (

Si)

+ η

80

2e 5

η +

δε

80

0p

6

L+

β γ 1

+ χ

80

0U

2β

+ γ 1

δ +

χ>

785

U3

β δ

+ χ

78

5e 6

L

(S

i) +

τ1

?e 7

max

L+

δ +

γ 1

ε7

72

P3

L

η +

τ2

~7

70

e 8m

ax

L+

δ ε

+ η

76

3U

4

η +

δ ε

?p

7m

in

Lτ 1

+ M

g2S

i +

(S

i)?

E1

L+

γ+τ 2

L+τ

1+τ

2γ+

τ 1+

τ 2

14

9

(Cu

)+β+

χL

+(C

u)+

χ

3

4

5

L+

β+γ 1

L+

δ+γ 1

β+δ+

γ 1

L+

γ 1+

χβ+

γ 1+

χ

13

γ 1+

δ+χ

β+δ+

χ

2

11

61

21

0

L+

δ+ε

L+

γ 1+

ε

8

δ+γ 1

+ε

1

L+η

+εη+

δ+ε

7τ 1

+M

g2S

i+(S

i)

233


MSIT®

Cu–Mg–Si

Fig

. 1b:

C

u-M

g-Si

.Par

tial

rea

ctio

n s

chem

e, p

art

2

Cu

-Si

Mg

-Si

A-B

-CC

u-M

g-S

iC

u-M

g

L

(S

i) +

τ2

?e 1

0m

ax

L+

εη

+ γ 1

74

5U

5

Lη

+ τ 2

+ (

Si)

73

9E

2

Lτ 1

+ τ

2+

(S

i)7

31

E3

δ +

χ γ 1

72

9p

8

l (

Cu

) + γ

72

5e 1

1

L

η +

γ~

723.5

e 12m

ax

Lγ

+ η

+ τ 2

72

3E

4

L +

(C

u)

χ +

γ7

18

U6

L

γ 1 +

γ7

11

e 13m

ax

δ ε

+ γ 1

71

0e 1

4

Lη

+ γ

+ γ 1

70

1E

5

Lχ

+ γ

+ γ 1

69

3E

6

χ +

γ (

Cu

) + γ

16

09

U7

l (

Mg)

+ M

g2S

i

637.6

e 15

L

Cu

Mg

2+τ

1

56

5e 1

6m

ax

L+

τ 1γ

+ C

uM

g2

552<

T<

56

5U

8

χ (

Cu

) + γ

1

55

2e 1

8

l C

uM

g2 +

γ5

52

e 17

L +

τ1

Cu

Mg

2 +

Mg

2S

i5

08

U9

l (

Mg)

+ C

uM

g2

48

5e 1

9

L

(M

g)

+ C

uM

g2 +

Mg

2S

i4

79

E7

(Mg)

+ C

uM

g2 +

Mg

2S

i

L

(C

u) +

γ~

760

e 9m

ax

η+τ 2

+(S

i)

ε+η+

γ 1L

+η+

γ 1

τ 1+τ

2+(

Si)

γ+η+

τ 2

L+χ

+γ(C

u)

+ χ

+ γ

η+γ+

γ 1

χ+γ+

γ 1

(Cu

)+χ+

γ 1(C

u)+

γ+γ 1

Cu

Mg

2+γ

+τ1

L+

CuM

g2+γ

τ 1+

CuM

g2+

Mg

2S

i

L+

CuM

g2+

Mg

2S

i

11

09

78

14

11

23

45

61

21

3

234


MSIT®

Cu–Mg–Si

70

80

10 20

20

30

Cu 90.00

Mg 0.00

Si 10.00

Cu 65.00

Mg 25.00

Si 10.00

Cu 65.00

Mg 0.00


Axes: at.%

(Si)e

5

U5

e3max

E2

e8max

η

800

750

900

e12max

τ2

E4

E5

E6

P3

e4

p4

p2 e

13max

U6

750

800

(Cu)γ

γ1

χ

δ

750

εP

2

U2

U4

β

U1

20

40

60

80

20 40 60 80

20

40

60

80

Cu Mg


Axes: at.%

e5

e9max

E7

e16max

η

800

750

1000

e 10max

τ1

E4

p2

1400

(Cu)γ

γ1

p1max

P1

940

1000

900

see Fig. 3

E2

E3

τ2

800

800

800

900

e19U

9U

8

e2max

Mg2Si

e7max

900

930

e17

e1

E1

e15

(Mg)CuMg2

800

800

900

1000

1100

1200

(Si)

900

1000

1100

1300

1200

1300

e4p

4

e11

Fig. 3: Cu-Mg-Si.

Liquidus surface in

the Cu corner

Fig. 2: Cu-Mg-Si.

Liquidus surface

235


MSIT®

Cu–Mg–Si

50

60

70

20 30 40

20

30

40

Cu 80.00

Mg 10.00

Si 10.00

Cu 40.00

Mg 50.00

Si 10.00

Cu 40.00

Mg 10.00


Axes: at.%

τ1+τ

2+τ

3

τ3

τ1

τ2

γ

τ1 + γτ

2+γ

τ2+τ

3

20

40

60

80

20 40 60 80

20

40

60

80

Cu Mg


Axes: at.%

Mg2Si

τ1+CuMg

2+Mg

2Si

CuMg2

CuMg2+Mg

2Si+(Mg)γ

ε

η"+τ2+(Si)

γ1

(Cu)

γ1+ε+γ

(Cu)+γ1+γ τ

3

η"

τ1+τ

2+(Si)

ε+η"+τ2

ε+τ2+γ

τ2

τ1

τ1+(Si)+Mg

2Si

τ1+τ

3+γ

τ1+Mg

2Si+γ

τ1+τ

2+τ

3

τ2+τ

3+γ

(Si)

Fig. 4: Cu-Mg-Si.


500°C in the 3 region

Fig. 5: Cu-Mg-Si.


450°C

236


MSIT®

Cu–Mg–Si

60 50 40 30

700

800

900

1000

Cu 66.70

Mg 33.30

Si 0.00

Cu 26.70

Mg 33.30

Si 40.00Cu, at.%

Te

mp

era

ture

, °C

L

γ

τ1+(Si)+Mg2Si

765

~950°C

transformationat 870-890

L+(Si)

L + τ

τ1

τ1+(Si)+L

p1, 930�H

,kJ

mo

l-1·

Mg

20 80706050403010 90

Mg, at.%

0

-5

-10

-20

-15

-25

-30

Cu

Si

Mg

70.00

30.00

0.00

Fig. 6: Cu-Mg-Si.

Partial vertical section

at 33.3 at.% Mg

Fig. 7: Cu-Mg-Si.

Experimental values

of the enthalpy of

mixing as a function

of composition for

liquid Cu-Mg-Si

alloys along the

isopleth with

xCu /xSi = 7/3;

dashed line is

interpolated part

237


MSIT®

Cu–Mg–Si

0

-1

-2

-4

-3

-5

0 20 40 1008060

Mg, at.%

lna

Mg

Fig. 8: Cu-Mg-Si.

Natural logarithm of

the magnesium

activity as a function

of composition for

liquid Cu-Mg-Si

alloys along the

isopleth with

xCu /xSi =7/3 at 900°C

238


MSIT®

Cu–Mg–Sn

Copper – Magnesium – Tin

Lazar Rokhlin and Evgeniya Lysova

Literature Data

Although a series of research groups have investigated the phase relations within the Cu-Mg-Sn ternary,

there is still no complete constitution diagram available. First information on the system was gathered in the

Cu-rich corner with respect to the possibility of age hardening [1927Dah, 1937Ven, 1938Per, 1940Per,

1948Ray] and were summarized in an early review by [1952Pie].

Detailed inspection of the crystal structures of ternary compounds [1950Kri, 1952Gla, 1955Gla, 1956Gla,

1960Now, 1976Osa, 1978Osa] and of the phase relations between 180 and 500°C [1955Tes, 1956Che,

1963Tes1, 1963Tes2] provided the knowledge for the construction of a partial isothermal section at 400°C

(<33 at.% Sn). The investigation of polythermal sections, Cu-CuMgSn, CuMg2-Mg2Sn, and

Cu2Mg-Mg2Sn, provided the basis for a partial liquidus projection and thus for a better understanding of

the solidification behavior [1963Tes1, 1963Tes2, 1969Phi]. These results were summarized in reviews by

[1979Dri, 1979Cha]. More recent research was devoted to refine the phase relations and established most

of the invariant reactions with high precision measurements [1991Vic, 1992Vic, 1996Vic].

[1927Dah] carried out thermal analysis, microstructure investigations and electrical conductance

measurements of Cu-rich alloys along the section with constant ratio Mg : Sn = 56.5 : 43.5 (at.%) containing

up to about 9.4 at.% Mg and 7.2 at.% Sn. The section constructed was considered incorrectly to be a part of

the pseudobinary system Cu-Mg2Sn with the eutectic point at about 12.2 at.% Mg and 9.4 at.% Sn and

690°C. The section furthermore showed a decrease of the limited joint solubility of Mg and Sn in solid

copper with decrease of temperature.

[1937Ven] studied alloys in the Cu corner of the Cu-Mg-Sn system and claimed Mg2Sn to be in equilibrium

with Cu. Excess of Mg, as compared with Mg2Sn, decreased the solubility of this compound in the Cu-base

solid solution, whereas small excess of Sn increased it slightly.

[1938Per, 1940Per] prepared Mg-rich Cu,Sn-alloys (up to 2.9 at.% Cu) by adding to Mg melts binary Cu-Sn

master alloys, Cu4Sn, Cu3Sn and Cu2Sn. The structure of the alloys, however, was not studied.

In a review of the work of [1927Dah], the existence of a pseudobinary section “Cu-Mg2Sn” was said to be

likely due to the high heat of formation of Mg2Sn as compared with the Cu-Sn compounds [1948Ray].

[1950Kri] studied the crystal structure of alloys along the section with 33.3 at.% Sn by X-ray diffraction

and discovered the ternary compound CuMgSn in coexistence with Mg2Sn, Cu2Mg and (Sn). Neither a

solid solution was formed with Mg2Sn nor dissolved Cu, Mg and Sn in CuMgSn to any appreciable

amounts. These results were later confirmed by [1960Now] from a study of the subsection CuMgSn -

Mg2Sn.

[1952Gla, 1955Gla] studied the alloys along the line Cu-CuMgSn by thermal analysis, X-ray diffraction and

microscopy and furthermore determined crystal structure and melting point of a new ternary congruently

melting compound Cu4MgSn.

[1955Tes] constructed the partial isothermal section of the Cu-Mg-Sn phase diagram at 400°C in the region

0-66.7 at.% Mg and 0-33.3 at.% Sn on a series of alloys annealed at 400°C employing optical microscopy

and X-ray diffraction methods. As a continuation, [1956Che] outlined in more detail the homogeneity

ranges of the solid phases (Cu), Cu2Mg and Cu4MgSn for 400°C and [1956Gla] confirmed the solid

solubility of Sn in Cu2Mg up to 15.0 at.% Sn at 400°C.

[1963Tes1] investigated the Cu-Mg-Sn phase diagram by microscopy, thermal analysis and X-ray

diffraction methods. The alloys were annealed at various temperatures within 180 to 500°C and cooled then

either by quenching or slowly down to room temperature. Three polythermal sections were constructed: 1)

Cu-CuMgSn, 2) CuMg2-Mg2Sn, 3) Cu2Mg-Mg2Sn, of which the polythermal sections Cu-CuMgSn and

CuMg2-Mg2Sn were recognized to be pseudobinary systems. Moreover, [1963Tes1] constructed an

isothermal section of the phase diagram in the solid state for the entire concentration range on the basis of

results obtained from alloys slowly cooled to room temperature.

239


MSIT®

Cu–Mg–Sn

[1963Tes1, 1963Tes2] carried out electrical conductivity measurements of the alloys belonging to the same

polythermal sections Cu-CuMgSn, CuMg2-Mg2Sn and Cu2Mg-Mg2Sn in order to redetermine the

boundaries of the Cu2Mg ternary solid solution and the homogeneity range of the ternary phase Cu4MgSn.

[1963Tes2] thereby confirmed earlier results by [1955Tes, 1956Che, 1956Gla, 1963Tes1]. The alloys were

annealed at 500°C for 10 days and then cooled to room temperature immediately in air or for 1 week.

[1969Phi] determined the Cu solid solution boundaries in the Cu-Mg-Sn system at 670°C by optical

microscopy and X-ray diffraction. Moreover, [1969Phi] determined the type and lattice parameter of the

cubic ternary phase in equilibrium with the Cu solid solution. The phase composition, however, remained

unknown.

[1976Osa] studied by scanning and transmission electron microscopy, optical microscopy and X-ray

diffraction the decomposition of the Cu supersaturated solid solution in a Cu-rich alloy with 2.1 at.% Mg

and 2.5 at.% Sn after homogenization at 700°C. At the late stages of the decomposition, precipitates of the

ternary phase Cu4MgSn occurred [1976Osa].

[1978Osa] refined the crystal structures of CuMgSn and of Cu4MgSn, discovered earlier by [1950Kri,

1952Gla].

Several reviews are due to [1952Pie, 1979Dri, 1979Cha]. Combining results of the previous investigations

[1979Cha] constructed a tentative isothermal section at 400°C and a projection of the liquidus surface of

the Cu-Mg-Sn phase diagram.

[1979Dri] generalized results of the investigations [1927Dah, 1950Kri, 1952Gla, 1963Tes1, 1963Tes2,

1956Che] and presented isothermal section of the Cu-Mg-Sn phase diagram at room temperature in whole

concentration range.

[1991Vic, 1992Vic] determined experimentally temperatures and compositions of two new eutectic

invariant equilibria in the Cu-Mg-Sn system, L (Mg)+CuMg2+Mg2Sn and L CuMg2+Mg2Sn. The quite

accurate study was performed by using differential thermal analysis, optical microscopy, electron probe

microanalysis and X-ray diffraction. Similarly, [1996Vic] established experimentally compositions of the

ternary eutectic point and the solid phases participating in the invariant equilibrium

L (Cu2Mg)+CuMg2+Mg2Sn suggested by [1979Cha]. Moreover, [1996Vic] proposed a partial liquidus

projection for the Cu2Mg-CuMg2-Mg2Sn-CuMgSn region using the binary phase diagrams of the Cu-Mg

and Mg-Sn systems, and the invariant equilibria L CuMg2+Mg2Sn and L (Cu2Mg)+CuMg2+Mg2Sn.

Binary Systems

The binary phase diagrams of the systems Cu-Mg, Mg-Sn are accepted from [Mas2], whilst Cu-Sn is from

[1990Sau].

Solid Phases

The solid phases established in the boundary binary systems and those involved in the studied parts of the

ternary Cu-Mg-Sn phase diagram are listed in Table 1. Copper dissolves in the adjoining binary systems up

to 6.93 at.% Mg and 9.1 at.% Sn at 725 and 596°C, respectively, [Mas2]. Magnesium dissolves in the

adjoining binary systems up to 0.013 at.% Cu and 3.35 at.% Sn at 485 and 561°C, respectively, [Mas2].

There is practically no solubility of Cu and Mg in solid tin ( Sn) [Mas2]. No homogeneity range is

established for the binary compound CuMg2 in the binary system Cu-Mg [Mas2], but in the ternary system,

electron microprobe analysis shows solubility of 2.2 at.% Sn in CuMg2 [1996Vic]. For Cu2Mg a significant

homogeneity range was established up to a maximum solubility of 15 at.% Sn at 400°C in direction to

CuMgSn [1955Tes, 1956Che, 1956Gla], 12 at.% Sn in direction to the Sn corner as reported by [1955Tes,

1956Che] and 2 at.% Mg at 552°C, 2.3 at.% Cu at 725°C in the binary system Cu-Mg [Mas2]. The

homogeneity range for Cu2Mg was determined to be 14.7 at.% Sn at 500°C and 13 at.% Sn at room

temperature [1963Tes2]. These solubility values agree, in general, with that reported by [1955Tes,

1956Che, 1956Gla] for 400°C. The limits of x and y values in the general formula Cu2+x-yMg1-xSny

presented in Table 1 for the Cu2Mg base compound correspond to the maximum homogeneity range in the

binary Cu-Mg system, to the maximum solubility of Sn at 400°C [1955Tes, 1956Che, 1956Gla] and to the

maximum width of the homogeneity area at 400°C in the ternary shown by [1956Che]. In the binary Cu-Sn

240


MSIT®

Cu–Mg–Sn

system all intermediate phases ( , 1, , , , , ') are characterized by certain homogeneity ranges whilst

Mg2Sn shows no homogeneity range in the binary [Mas2].

Mg2Sn, according to [1996Vic], dissolves up to about 1 at.% Cu. The homogeneity range of the ternary

compound CuMgSn was not established. The concentration limits of the large homogeneity range of

Cu4MgSn are shown in Table 1 for 400°C according to [1955Tes]. The homogeneity range of Cu4MgSn

was also shown by [1956Che], as result of an investigation conducted with the same details as [1955Tes].

Data of [1955Tes] are preferable, however, as compared with [1956Che], because they are more compatible

with the polythermal section Cu-CuMgSn studied by [1963Tes1, 1963Tes2].


Two pseudobinary sections are established reliably in the Cu-Mg-Sn phase diagram: CuMg2-Mg2Sn and

Cu-Cu4MgSn ( 2). The section CuMg2-Mg2Sn is shown in Fig. 1. It is drawn mainly after [1963Tes1], but

with some corrections to meet the melting temperatures of the compounds CuMg2 and Mg2Sn according to

the binary systems [Mas2] and temperature and composition of the eutectic point accepted from [1991Vic].

Temperature and composition of the eutectic point are determined in [1991Vic] with good precision. The

second pseudobinary section Cu-Cu4MgSn ( 2) is shown in Fig. 2 as part of the partial polythermal section

running from Cu to CuMgSn ( 1). It is a reproduction from [1963Tes1] because continuation of the

pseudobinary section Cu-Cu4MgSn ( 2) beyond Cu4MgSn ( 2) may be considered conditionally to be also

a pseudobinary section up to point p4 corresponding to the liquid involved in the peritectic formation of

CuMgSn ( 1). The part of the section from Cu4MgSn ( 2) to the point p4 may be considered to be

pseudobinary only conditionally because it is not limited by two congruent melting compounds, which

could be assumed as components. However, the peritectic three-phase equilibrium corresponding to the

CuMgSn formation can be considered to be invariant.


Two four-phase and three three-phase invariant equilibria were established in the Cu-Mg-Sn system and are

listed in Table 2. Figure 3 shows the partial reaction scheme. Characteristics of the three-phase invariant

equilibria e2(max) and p4 are assumed after [1963Tes1]. Characteristics of the other invariant equilibria

e5(max), E1, E2, are assumed after [1991Vic, 1996Vic] where they were determined most exactly. The

compositions of the solid phases , CuMg2 and Mg2Sn participating in the equilibrium E1 are accepted from

[1996Vic] determined by electron probe microanalysis. There is no special determination of the CuMg2 and

Mg2Sn compositions in the equilibria e5(max) and E2, but they are assumed to correspond to those in E1 in

[1996Vic]. The composition of the magnesium solid solution (Mg) in E2 is extracted from the binary

systems Cu-Mg and Mg-Sn at the equilibrium temperature [Mas2].

Liquidus Surface

The projection of the liquidus surface of the Cu-Mg-Sn phase diagram is shown in Fig. 4. It is essentially

follows [1979Cha] with some corrections of the eutectic points due to [1991Vic, 1996Vic]. The liquidus

surface is based on insufficient experimental data and at present is considered to be tentative. Therefore, all

monovariant equilibrium lines are shown by dashed lines. Moreover, some border lines to the fields of

primary phase crystallization are missed.

Isothermal Sections

Figure 5 displays the isothermal section of the Cu-Mg-Sn phase diagram at 400°C. It is mainly from

[1955Tes] with some corrections to meet the accepted binary phase diagrams and the homogeneity ranges

of the phases CuMg2 and Mg2Sn established by [1996Vic]. Moreover, the section contains some additions

made by [1979Cha]. The section misses, however, a number of the phase fields for the binary Cu-Sn phases,

, and . Necessary existence of these fields arise from the binary Cu-Sn system [Mas2]. Due to lack of

experimental data these phase fields can not be shown.

241


MSIT®

Cu–Mg–Sn


Figure 6 presents the vertical section Cu2Mg-Mg2Sn of [1963Tes1] taking into account the accepted binary

phase diagrams [Mas2] and the established homogeneity ranges of the solid phases. The section crosses the

eutectic valley L +Mg2Sn. The eutectic character of this monovariant equilibrium was stated by

[1963Tes1] and confirmed by microstructural investigation [1996Vic].

Thermodynamics

[1972Pre] determined enthalpies of formation of the Cu2Mg type phase in the ternary system Cu-Mg-Sn

using calorimetry and liquid tin as solvent. The study showed an increase of the phase stability replacing

Cu atoms successively for Sn atoms.

[1987Sir, 1995Som] studied the ability of Cu-Mg-Sn ternary alloys to form glasses in high speed

solidification.

References

[1927Dah] Dahl, O., “About the Structure and Age Hardening of the Cu-Rich Cu-Mg and Cu-Mg-Sn

Alloys” (in German), Wiss. Veroeff. Siemens-Konzern, 6, 222-234 (1927/1928) (Equi.

Diagram, Experimental, *, 3)

[1937Ven] Venturello, G., Fornaseri, M., “Magnesium-Tin-Copper Alloys Rich in Copper” (in Italian),

Metall. Ital., 29(5), 213-221 (1937) (Equi. Diagram, Experimental)

[1938Per] Peredelskii, K.V., “Magnesium Alloys with Intermetallic Compounds” (in Russian),

Aviapromyshlennost, (6), 24-26 (1938) (Experimental, 7)

[1940Per] Peredelskii, K.V., “About Alloys of the Intermetallic Compounds with Magnesium” (in

Russian), Metallurg, 15(6), 30-34 (1940) (Experimental, 11)

[1948Ray] Raynor, G.V., “Equilibrium Relationships in Ternary Alloys”, Philos. Mag., 39(290),

218-229 (1948) (Review, 12)

[1950Kri] Kripyakevich, P.I., Gladyshevsky, E.I., Cherkashin, E.E., “The Crystal Structure of the

CuMgSn Ternary Phase” (in Russian), Dokl. Akad. Nauk SSSR, 75(2), 205-207 (1950)

(Crys. Structure, Experimental, *, 7)

[1952Gla] Gladyshevsky, E.I., Kripyakevich, P.I., Teslyuk, M.Yu., “The Crystal Structure of the

Cu4MgSn Ternary Phase” (in Russian), Dokl. Akad. Nauk SSSR., 85(1), 81-84 (1952) (Crys.

Structure, Experimental, *, 8)

[1952Pie] Pietch, E.H.E., Meyer, R.J., “Magnesium-Copper-Tin Alloys” (in German), Gmelins

Handbuch der Anorg. Chemie, Verlag Chemie, GmbH., Weinheim, Germany, 27(A4), 722

(1952) (Review, 3)

[1955Gla] Gladyshevsky E.I., Kripyakevich, P.I., “Position of the Cu and Mg Atoms in the Structure

of CuMgSn” (in Russian), Dokl. Akad. Nauk SSSR, 102(4), 743-746 (1955) (Crys.

Structure, 6)

[1955Tes] Teslyuk, M.Yu., Gladyshevsky, E.I., “Solubility of Tin in the Intermetallic Phase Cu2Mg”

(In Russian), Naukovi Zapisky L'viv. Derz. Univer., Ser. Khim., 34(4), 84-90 (1955) (Equi.

Diagram., Crystal Structure, Experimental, #, 13)

[1956Che] Cherkashin, E.E., Gladyshevsky, E.I., Teslyuk, M.Yu., “Investigation of the Cu-Mg-Sn

System in the Cu-Cu2Mg-CuMgSn Region” (in Russian), Izv. Sekt. Fiz.-Khim. Anal., Akad.

Nauk SSSR, 27, 212-216 (1956) (Crys. Structure, Experimental, Theory, *, 11)

[1956Gla] Gladyshevsky, E.I., Cherkashin, E.E., “Solid Solutions Based on Metallic Compounds” (in

Russian), Russ. J. Inorg. Chem., 1(6), 288-295 (1956), translated from Zh. Neorg. Khimii,

1(6), 1394-1401 (1956) (Review, 4)

[1960Now] Nowotny, H., Holub, F., “Investigation of Metallic Systems with Fluorspar Phases” (in

German), Monatsh. Chem., 91(5), 877-887 (1960) (Crys. Structure, Review, *, 15)

[1963Tes1] Teslyuk, M.Yu., “A Study of the Magnesium-Copper-Tin Ternary System” (in Ukrainian),

Zb. Rob. Asp. Lv. Uni. Pri. Nauk, Lvov, 5-10 (1963) (Equi. Diagram, Experimental, #, 10)

242


MSIT®

Cu–Mg–Sn

[1963Tes2] Teslyuk, M.Yu., “Electrical Conductivity Study of Certain Sections of Cu-Mg-Sn” (in

Ukrainian), Zb. Rob. Asp. Lv. Uni. Pri. Nauk, Lvov, 11-14 (1963) (Experimental, *, 9)

[1969Phi] Phillips, D.L., Ainsworth, P.A., “Age Hardenable Sn-Mg Bronze” (in German), Metall,

23(8), 804-808 (1969)

[1972Pre] Predel, B., Ruge, H., “Study of the Enthalpies of Formation in the Mg-Cu-Zn, Mg-Cu-Al

and Mg-Cu-Sn Systems as a Contribution to the Understanding of the Binding Conditions

of Laves Phases” (in German), Mater. Sci. Eng., 9(3), 141-151 (1972) (Experimental,

Thermodyn., Review, 61)

[1976Osa] Osamura, K., Takamuku, S., Murakami, Y., “The Structure and Behavior of Precipitates in

a Cu-Sn-Mg Ternary Alloy” (in German), Z. Metallkd., 67(7), 467-472 (1976)


[1978Osa] Osamura, K., Murakami, Y., “Crystal Structure of CuSnMg and Cu4SnMg Ternary

Compounds”, J. Less-Common Met., 60(2), 311-313 (1978) (Experimental, Crys.

Structure, *, 6)

[1979Cha] Chang, Y.A, Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Mg-Sn” in “Phase Diagram

and Thermodynamic Properties of Ternary Copper-Metal Systems”, INCRA Monograph

Series 6,520-525 (1979) (Equi. Diagram, Review, #, 9)

[1979Dri] Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L.,

Turkina, N.I., “Copper-Magnesium-Tin” (in Russian), in “Binary and Multicomponent

Copper-Base Systems”, Nauka, Moskow, 163-164 (1979) (Review, 7)

[1984Nay] Nayeb-Hashimi, A. A., Clark, J. B., “The Cu-Mg (Copper-Magnesium) System”, Bull.

Alloy Phase Diagrams, 5(1), 36-43 (1984) (Review, Equi. Diagram, 42)

[1987Sir] Sirkin, H., Mingolo, N., Nassif, E., Arcondo, B., “Increase of the Glass-Forming

Composition Range of Mg-Based Binary Alloys by Addition of Tin”, J. Non-Cryst. Solids,


[1990Sau] Saunders, N., Miodownik, A. P., “The Cu-Sn (Copper-Tin) System”, Bull. Alloy Phase

Diagrams, 11(3), 278-287 (1990) (Equi. Diagram, Review, 57)

[1991Vic] Vicente, E.E., Bermudez, S., Esteban, A., Tendler, R., Arcondo, B., Sirkin, H., “Invariant

Three- and Four-Phase Equilibria in the Magnesium-Rich Corner of the Mg-Cu-Sn Ternary

System”, J. Mater. Sci., 26(5),1327-1332 (1991) (Equi. Diagram, Experimental, #, 12)

[1992Vic] Vicente, E.E., Bermudeez, S., Tendler, R., Arcondo, B., Sirkin, H., “The

Mg-Mg2Cu-Mg2Sn Ternary Eutectic”, J. Mater. Sci. Lett., 11(8), 523-524 (1992) (Equi.


[1995Som] Somoza, J.A., Gallego, L.J, Rey, C., Rosenberg, S., Arcondo, B., Sirkin, H., De Tendler,

R.H., Kovacs, J.A., Alonso, J.A., “An Experimental and Theoretical Study of the

Glass-Forming Region of the Mg-Cu-Sn System”, J. Mater. Sci., 30, 40-46 (1995)

(Experimental, Theory, 37)

[1996Vic] Vicente, E.E., Bermudez, S., De Tendler, R.H., Arcondo, B., Sirkin, H., “The

Cu2Mg-CuMg2-Mg2Sn Ternary Eutectic”, J. Mater. Sci. Lett., 15(19), 1690-1696 (1996)

(Equi. Diagram, Experimental, #, 15)

243


MSIT®

Cu–Mg–Sn


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.94

c = 521.07

pure Mg at 25°C [Mas2]

(Cu)

< 1084.87

cF4

Fm3m

Cu

a = 361.46 pure Cu at 25°C [Mas2]

( Sn)

231.9681 - 13

tI4

I41/amd

Sn

a = 583.18

c = 318.18

pure Sn at 25°C [Mas2]

( Sn)

< 13

cF8

Fd3m

C (diamond)

a = 648.92 [Mas2]

CuMg2

< 568

oF48

Fddd

CuMg2

a = 905

b = 1824.7

c = 528.3

[Mas2, 1984Nay]

, Cu2+x-yMg1-xSny

Cu2Mg

< 797

cF24

Fd3m

Cu2Mg

a = 703.5

a = 726.3

-0.06 x 0.13,

0 y 0.45 [Mas2, 1956Che]

at x = 0, y = 0 [Mas2, 1984Nay]

[1955Tes, 1956Che]

from kX-units

x = 0, y = 0.45 (Cu51.7Mg33.3Sn15)

linear da/dx

, Cu-Sn

798-586

cI2

Im3m

W

a = 297.81 to

298.71

13.1-16.5 at.% Sn [Mas2, 1990Sau]

1, Cu-Sn

755 - 520

cF16

Fm3m

BiF3

a = 606.05 to

611.76

15.5-27.5 at.% Sn [Mas2, 1990Sau]

, Cu-Sn

590 - 350

cF416

F43m

Cu41Sn11

a = 1798.0 20-21 at.% Sn [Mas2, 1990Sau]

, Cu-Sn

640 - 582

hP26

P63/m

Cu10Sn3

a = 733.0

c = 786.4

20.3-22.5 at.% Sn [Mas2, 1990Sau]

, Cu-Sn

< 676

oC80

Cmcm

Cu3Sn

a = 552.9

b = 4775.0

c = 432.3

24.5-25.9 at.% Sn [Mas2, 1990Sau]

, Cu-Sn

415 - 186

hP4

P63/mmc

NiAs

a = 419.0

c = 508.6

43.5-45.5 at.% Sn [Mas2, 1990Sau]

244


MSIT®

Cu–Mg–Sn


', Cu-Sn

< 189

h**

ordered

NiAs-deriv.

a = 2087.0

c = 2508.1

45 at.% Sn [Mas2, 1990Sau]

Mg2Sn

< 770.5

cF12

Fm3m

CaF2

a = 676.38 [Mas2, 1984Nay]

* 1, CuMgSn

< 610

cF12

F43m

AgMgAs

a = 622.4

a = 626.2

a = 6.226

a = 6.228

[1963Tes1, 1978Osa]

[1950Kri]

[1955Gla] from kX

[1960Now] from kX

* 2,

Cu4-x+yMg1+xSn1-y

< 750

Cu4MgSn

cF24

Fd3m

Cu2Mg-derivative

a = 704.2

0 x 0.50,

-0.05 y 0.25 [1955Tes]

at x = 0, y = 0 [1978Osa]


Cu Mg Sn

L (Cu) + 2 720 e2(max) L

(Cu)

2

76.5

91.7

69.2

11.75

4.15

15.4

11.75

4.15

15.4

L + 2 1 610 p4 L

2

1

31.8

65.2

33.4

34.1

17.4

33.3

34.1

17.4

33.3

L CuMg2 + Mg2Sn 522 e5(max) L

CuMg2

Mg2Sn

66.3

32.9

1.0

26.0

64.9

64.9

7.7

2.2

34.1

L + CuMg2 + Mg2Sn 520 E1 L

CuMg2

Mg2Sn

29.8

50.5

32.9

1.0

62.4

42.4

64.9

64.9

7.8

7.1

2.2

34.1

L CuMg2 + Mg2Sn + (Mg) 467 E2 L

CuMg2

Mg2Sn

(Mg)

82.1

32.9

1.0

0.01

13.5

64.9

64.9

98.2

4.4

2.2

34.1

1.8

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

245


MSIT®

Cu–Mg–Sn

10 20 30

400

500

600

700

800

Cu 33.30

Mg 66.70

Sn 0.00

Cu 0.00

Mg 66.70

Sn 33.30Sn, at.%

Te

mp

era

ture

, °C

522

CuMg2 + Mg2Sn

Mg2Sn + L

L

770.5°C

568°C

e5(max)

CuMg2+L

Fig. 1: Cu-Mg-Sn.

The pseudobinary

system

CuMg2-Mg2Sn

10 20 30

500

600

700

800

900

1000

1100

Cu Cu 24.00

Sn 38.00

Mg 38.00Sn, at.%

Te

mp

era

ture

, °C

610

1084.87°C

720

L

τ1+L

4.15

e2(max)

(Cu)

(Cu) + τ2

750

τ2

τ2 + τ1τ1

Fig. 2: Cu-Mg-Sn.

The pseudobinary

system Cu-CuMgSn

246


MSIT®

Cu–Mg–Sn

Fig. 3: Cu-Mg-Sn. Partial reaction scheme

Cu-Mg A-B-CCu-Mg-Sn Mg-Sn

l CuMg2 + γ

552 e4

l (Mg)+CuMg2

485 e6

L CuMg2 + τ

2

720 e2max

L + τ2

τ1

610 p4

L CuMg2+Mg

2Sn

522 e5max

L γ+CuMg2+Mg

2Sn520 E

1

L CuMg2+Mg

2Sn+(Mg)467 E

2

l (Mg)+Mg2Sn

561.2 e3

L+γ+Mg2Sn

γ+CuMg2+Mg

2Sn

CuMg2+Mg

2Sn+(Mg)

20

40

60

80

20 40 60 80

20

40

60

80

Cu Mg

Sn Data / Grid: at.%

Axes: at.%

p1

e7

e8

(Sn)

η

ε

β

p4

600

1000

p5

600p

3

p2

τ1

e1

γ1

γτ

2

e2

(Cu)

800

700

900

Mg2Sn

700

e3E

2

Mg2Sn

e6

E1

e5

e4

600

600

600

(Mg)CuMg2

Fig. 4: Cu-Mg-Sn.

Projection liquidus

surface

247


MSIT®

Cu–Mg–Sn

10 20 30

400

500

600

700

800

Sn 0.00

Cu 66.70

Mg 33.30

Sn 33.30

Cu 0.00

Mg 66.70Sn, at.%

Te

mp

era

ture

, °C

620

Mg2Sn+L

L

797°C

γ + L

770.5°C

520

18.5

γ+Mg2Sn+L

γ+CuMg2

~8

~9.25

γ+CuMg2+Mg2Sn

γγ+CuMg2+L

20

40

60

80

20 40 60 80

20

40

60

80

Cu Mg


Axes: at.%

η

εδ

τ1

γ+CuMg2+Mg

2Sn

(Cu)+τ2+γ

Cu2Mg(Cu)

Mg2Sn

CuMg2

CuMg2+Mg

2Sn+(Mg)

L

γ

γ+τ1+Mg

2Sn

τ2

τ 2+γ+τ 1

(Mg)

Fig. 6: Cu-Mg-Sn.

Vertical section

Cu2Mg-Mg2Sn

Fig. 5: Cu-Mg-Sn.


400°C

248


MSIT®

Cu–Mg–Ti

Copper – Magnesium – Titanium

Lazar Rokhlin and Nataliya Bochvar

Literature Data

[1966Zwi] studied Cu-rich alloys of the Cu-Mg-Ti system containing 0-22 at.% Ti and 0-36 at.% Mg.

Conditions for preparation of the alloys were described earlier [1965Zwi]. They included use of electrolytic

Cu of common purity, pure Mg and master alloy of Cu with 42 mass% Ti. Melting of the alloys were carried

out in an arc furnace under Ar atmosphere or in a Tammann furnace in graphite crucibles followed by

casting the melt into steel moulds. The cast alloys were then sealed in quartz ampoules and annealed at

700°C for 160 h and quenched in water. [1966Zwi] did not report on the methods of investigation: most

likely, thermal analysis, light optical microscopy and electron probe microanalysis were used. As a result

of the investigation, [1966Zwi] constructed the liquidus surface for the studied concentration area and a

partial isothermal section of the phase diagram at 700°C. Furthermore [1966Zwi] established the formula

“Ti2Cu7” for the binary compound richest in Cu in contrast to [Mas2], who assumed for this compound the

formula TiCu4.

The reviews [1979Dri, 1981Dri] include descriptions of the Cu-Mg-Ti phase diagram after [1966Zwi].

Binary Systems

The binary systems Cu-Mg and Cu-Ti are assumed according to [Mas2]. The binary system Mg-Ti is

beyond the studied concentration area of the ternary system.

Solid Phases

No ternary compound is found in the studied part of the ternary system Cu-Mg-Ti. Characteristics of the

relevant binary compounds Cu2Mg and TiCu4 are given in Table 1. The solid solubility of Ti in Cu2Mg and

the solid solubility of Mg in TiCu4 are not reported. Ti and Mg slightly increase their mutual solubility in

solid Cu [1966Zwi].


Only one four-phase invariant equilibrium of transition type is established in the studied part of the system.

Characteristics of the invariant equilibrium are given in Table 2 in accordance with [1966Zwi].

Concentrations of the liquid phase and copper solid solution (Cu) taking part in the invariant equilibrium

are essentially accepted from the drawings presented by [1966Zwi] with slight modification with respect to

the accepted binary systems. The respective partial reaction scheme is presented in Fig. 1.

Liquidus Surface

The projection of the partial liquidus surface for the Cu corner of the Cu-Mg-Ti phase diagram is presented

in Fig. 2 including lines of double saturation and the invariant equilibrium plane.

Isothermal Sections

Figure 3 displays the partial isothermal section of the Cu-Mg-Ti phase diagram constructed by [1966Zwi]

with correction of the Cu2Mg and TiCu4 homogeneity ranges to comply with the binary systems Cu-Mg

and Cu-Ti accepted from [Mas2].

Miscellaneous

[1965Zwi] reported results of the hardness and conductance measurements of the ternary alloys Cu-Mg-Ti

with about 2 mass% Ti and 1-2 mass% Mg. Measurements were carried out at room temperature after

249


MSIT®

Cu–Mg–Ti

ageing at 350 to 550°C for 30 to 10000 min. In addition to this, oxidation of the alloys during long-time

annealing at 700°C in air was studied.

References

[1965Zwi] Zwicker, U., Vogl, H., “On the Effect of Magnesium on Ageing and Oxide Scale Formation

of Copper-Titanium Alloys” (in German), Metall, 19(11), 1173-1178 (1965)

(Experimental, 16)

[1966Zwi] Zwicker, U., Kalsch, E., Nishimura, T., Ott, D., Seilstorfer, H., “Effect of Additions on

Equilibria in Copper-Rich Copper-Titanium Alloys” (in German), Metall, 20(12),


[1979Dri] Drits, M.E., Bochvar, N.R., Guzei, L.S., Lusova, E.V., Padezhnova, E.M., Rokhlin, L.L.,

Turkina, N.I., “Copper-Magnesium-Titanium”, in “Binary and Multicomponent

Copper-Base Systems” (in Russian), Moscow, Nauka, 166-167 (1979) (Equi. Diagram,

Review, 1)

[1981Dri] Drits, M.E., Bochvar, N.R., Guzei, L.S., Lusova, E.V., Padezhnova, E.M., Rokhlin, L.L.,

Turkina, N.I., “The Cu-Mg-Ti System (Copper - Magnesium - Titanium)”, Bull. Alloy

Phase Diagrams, 2(2), 226 (1981) (Equi. Diagram, Review, 1)



Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comment/References

(Cu)

< 1084.87

cF4

Fm3m

Cu

a = 361.48 pure Cu at 25°C

[Mas2, V-C2]

Cu2Mg

< 797

cF24

Fd3m

Cu2Mg

a = 701.3 [Mas2, V-C2]

TiCu4

< 885 - 400

oP20

Pnma

ZrAu4

a = 454.5

b = 434.1

c = 1295.3

[Mas2, V-C]

TiCu4

500

tI10

I4/m

MoNi4

a = 584.0

c = 362.0

[Mas2, V-C]


Cu Mg Ti

L + TiCu4 (Cu) + Cu2Mg ~780 U1 L

TiCu4

(Cu)

Cu2Mg

~76

80.9

89.6

66.7

~23

-

6.7

33.3

~1

19.1

3.7

-

250


MSIT®

Cu–Mg–Ti

Fig. 1: Cu-Mg-Ti. Partial reaction scheme

Cu-Mg A-B-CCu-Mg-Ti Cu-Ti

l (Cu)+Cu2Mg

e1

725

l+(Cu) βTiCu4

p1

885

L+βTiCu4

(Cu)+Cu2Mg780 U

1

βTiCu4+(Cu)+Cu

2Mg

L+βTiCu4+Cu

2Mg

L+(Cu)+Cu2Mg

10

20

30

40

60 70 80 90

10

20

30

40

Ti 50.00

Cu 50.00

Mg 0.00

Cu

Ti 0.00

Cu 50.00

Mg 50.00Data / Grid: at.%

Axes: at.%

βTiCu4

(Cu)

e1

U1

p1

Cu2Mg

βTiCu4

Cu2Mg

Fig. 2: Cu-Mg-Ti.

Partial liquidus

projection for

four-phase

equilibrium and edges

of double saturation

251


MSIT®

Cu–Mg–Ti

10

20

30

40

60 70 80 90

10

20

30

40

Ti 50.00

Cu 50.00

Mg 0.00

Cu

Ti 0.00

Cu 50.00

Mg 50.00Data / Grid: at.%

Axes: at.%

βTiCu4

βTiCu4+(Cu)

(Cu)

Cu2Mg

(Cu)+Cu2Mg

βTiCu4+Cu

2Mg

(Cu)+βTiCu4+Cu

2Mg

Fig. 3: Cu-Mg-Ti.

Partial isothermal

section at 700°C after

[1966Zwi]

252


MSIT®

Cu–Mg–Zn

Copper – Magnesium – Zinc

Hans Leo Lukas, Peter Rogl, Gautam Ghosh, Günter Effenberg

Literature Data

In the earliest study [1906Gui] established the structure-property relationship of Cu-40 mass% Zn-x mass%

Mg (up to 0.35 mass%) alloys. The first comprehensive study of the ternary phase equilibria was carried out

by [1948Koe]. They used electrolytic Cu (unspecified purity), 99.99% purity Zn and pure magnesium

(unspecified purity) to prepare the ternary alloys. [1948Koe] determined the liquidus surface and two

pseudo-binary sections by metallography and thermal analysis. [1950Mik] also reported the liquidus

surface. In addition [1950Mik] extensively studied solid phase equilibria and presented their results in

several vertical sections. The pseudobinary section Cu2Mg-MgZn2 reported by [1948Koe] was later

confirmed by [1952Lie]. The latter authors prepared 17 ternary alloys along the Cu2Mg-MgZn2 section

using 99.9% pure Mg and Cu, pure Zn (unspecified purity). The pseudobinary section and the phase

characterization were performed using metallography, thermal analysis and X-ray diffraction techniques.

[1955Gla, 1956Gla1, 1956Gla2] reported a partial isothermal section at 400°C. [1972Yam] investigated the

Zn corner using 52 ternary alloys containing 0.6 to 32.2 mass% Cu and 0.1 to 13.6 mass% Mg. The ternary

alloys were prepared using 99.998% purity Cu and Zn and 99.9% purity Mg. The liquidus surface of the Zn

corner and several vertical sections were determined by means of thermal analysis and metallography

[1960Wat, 1972Yam]. Most of these results were reviewed by [1979Cha]. Thermodynamic datasets of the

ternary system were assessed by [1997Lia] and [1998Lia]. [1998Lia] determined also the Cu-solubilities of

the MgZn, Mg2Zn3 and Mg2Zn11 phases by EDX.

Binary Systems

The Cu-Mg binary phase diagram is accepted from [1984Nay]. The Cu-Zn and Mg-Zn binary phase

diagrams are accepted from [Mas2]. Thermodynamically assessed datasets of all three binary systems are

accepted from the COST 507 action [1998Ans]. They agree well with the phase diagrams of [1984Nay] and

[Mas2], respectively; the temperatures of invariant equilibria do not differ more than 3 K.

Solid Phases

Structural data of all solid phases are given in Table 1. All three types of Laves phases [1934Lav] exist in

this ternary system: C14 (MgZn2), C15 (Cu2Mg), and C36 ( or CuMg2Zn3). The relative stability of these

three phases is influenced by electronic factors [1936Lav, 1970Kom]. Magnetic susceptibility [1954Kle,

1958Moe] and hydrogen solubility [1954Lie, 1957Sie] measurements of alloys along the Cu2Mg-MgZn2

section corroborate the importance of electronic factors and indicate variations in the density of states which

correlate with the composition ranges of stability. The lattice parameter of the C15 ((Cu1-xZnx)2Mg) Laves

phase has been measured several times [1952Lie, 1955Gla, 1971Sha, 1979Ell]. The effect of replacing Cu

atoms by Zn atoms on this lattice parameter is shown in Fig. 1. The lattice parameter values of [1952Lie]

are significantly lower than those of [1955Gla], [1971Sha], [1979Ell]. The best fit straight line for the latter

three sets of data can be expressed as a (in pm)=703.40+27.8x, referred to x of the formula (Cu1-xZnx)2Mg.

The stacking sequences of the C14, C15 and C36 Laves phases can be formulated as AB', ABC, and AB'A'C,

respectively. They are also called 2-layer (2H), 3-layer (3C) and 4-layer (4H) structures, respectively.

Between the homogeneity ranges of the C14 (MgZn2) and C36 phases, [1970Kom] reported the existence

of stacking variant structures with 8 layers (8H) (AB'AB'A'CA'C), 9 layers (9R) (AB'ABC'BCA'C) and 10

layers (10H) (ABC'BCA'C'BC'B'). These stacking variant structures can also be stabilized under pressure.

[1975Nak] studied the effect of pressure (up to 90 kbar) on the stability of (Cu1-xZnx)2Mg alloys at x = 0.8,

0.925 and 1. They found 9-layer (9R) 10-layer (10H) 4-layer (4H) phase transitions with increasing

pressure. Also, with increasing pressure both axes, c and a, decrease monotonically [1975Nak, 1977Oom].

Within the homogeneity range of the C15-phase there is no evidence of a structural transition [1979Ell].

The stacking variant structures of Laves phases in various systems have been discussed by [1977Ray].

253


MSIT®

Cu–Mg–Zn

At 400°C the Cu solubility in MgZn2 is reported to be about 3 at.% Cu [1948Koe, 1998Lia] or about 1 mol%

Cu2Mg. At 335°C the Cu solubilities in MgZn, Mg2Zn3 and Mg2Zn11 are about 2, 1 and 6 at.%, respectively

[1998Lia]. dissolves about 4 at.% Mg and ' dissolves about 2 at.% Mg [1955Gla].


The pseudobinary section Cu2Mg-MgZn2 is well established [1948Koe, 1952Lie]. It is shown in Fig. 2,

calculated using the dataset of [1998Lia]. The calculation reproduces the pseudobinary peritectics not

exactly in the plane of the section, but at 33.2 and 33.35 at.% Mg respectively, therefore in Fig. 2 they

deviate slightly from horizontal lines. The liquidus temperatures measured for Zn-contents less than 33 at.%

by [1950Mik] and [1952Lie] scatter significantly more than those of [1948Koe], which therefore are

accepted here for this composition range. The C15 (Cu2Mg) Laves phase melts congruently, [1948Koe]

localized the maximum at 850°C and about 33 at.% Zn by interpolating between their experimental points.

The calculation [1998Lia], which is in principle also an interpolation between the same experimental points,

indicates a slightly higher temperature (859°C) and lower Zn concentration (28 at.%) for the maximum. At

the equiatomic composition also the electric resistivity and its temperature dependence show maxima

[1956Mik]. However, [1956Mik] misinterpreted the alloy of equiatomic composition as a distinct ternary

phase. A ternary Laves phase (CuMg2Zn3) forms at a higher Zn-concentration of 53 at.% by a peritectic

reaction L+Cu2Mg at 718°C. Figure 3 shows another section, from Mg to Cu5Zn8, calculated using the

dataset of [1998Lia]. This section was reported to be pseudobinary by [1948Koe, 1950Mik], however, it

intersects the C15 (Cu2Mg) Laves phase somewhat off the congruent melting composition and therefore is

only approximately a pseudobinary one.


The reaction scheme for the solidification of ternary alloys is shown in Fig. 4, it is calculated using the

dataset of [1998Lia]. The invariant equilibria are also given in Table 2. The maximum e1 between e2 and

U1 is omitted in Fig. 4 for clarity. The reaction schemes given in the works of [1948Koe], [1949Mik] and

[1972Yam] differ in some details. [1948Koe] for simplification did not distinguish the different Laves

phases and neglected the participation of Mg2Zn3 and MgZn in the ternary equilibria. They gave the same

reactions for U1 (705°C), U4 (520°C) and U8 (370°C), U2 they described as eutectic at 700°C, E2 and P1 as

U type reactions at 452°C and 375°C, respectively. [1972Yam] found U8 at 367°C, but gave different

reaction sequences instead of E1+U4 and of P1+U6 leading to two three-phase equilibria +MgZn2+ and

Mg2Zn11+MgZn2+ at temperatures below 340°C, different from those published by [1948Koe] and

calculated from the dataset of [1998Lia], +MgZn2+Mg2Zn11 and + +Mg2Zn11.

Liquidus and Solvus Surfaces

Figure 5 shows the liquidus surface calculated from the dataset of [1998Lia]. An enlarged view of the

Zn-corner is shown in Fig. 6. Topologically the liquidus surface diagrams published by [1948Koe,

1960Wat] and [1972Yam] are equivalent, except [1948Koe] did not distinguish the different Laves phases

and in the diagram of [1972Yam] the field of primary crystallization of reaches that of and the sequence

P1 U6 p9 is replaced by opposite sequence of temperatures.

The dataset of [1998Lia] allows the calculation of the solidus and solvus surfaces of the (Cu) and (Zn) solid

solutions. They are equivalent to thermodynamic extrapolations from the binary subsystems as no detailed

measurements of ternary solubilities are published. The Cu solubility in (Mg) is not well known even in the

binary Cu-Mg system, therefore the (Mg)-solidus and -solvus cannot be calculated reliably. The solvus

surfaces of (Cu) and (Zn) are shown in Figs. 7 and 8, respectively.

Isothermal Sections

Figure 9 shows the isothermal section at 340°C, calculated from the dataset of [1998Lia]. This temperature

is just below the lowest liquidus temperature of the system (341°C at E3).

254


MSIT®

Cu–Mg–Zn


In addition to the liquidus surface and Mg-Cu5Zn8 pseudobinary section, [1950Mik] reported the following

vertical sections: Cu-MgZn, Zn-MgCu, Mg2Cu-CuZn2, Cu-MgZn2, at 10 at.% Mg, at 20 at.% Mg, at 30

at.% Mg, at 50 at.% Mg, at 60 at.% Mg, at 70 at.% Mg, and at 90 at.% Mg. [1960Wat] reported vertical

sections at 5 mass% Cu and 3 mass% Mg. [1972Yam] determined vertical sections at 1.7 mass% Cu, 6

mass% Cu, 15 mass% Cu, 2 mass% Mg, and 5 mass% Mg.

Thermodynamics

The standard enthalpy of formation ( Hf298) of Cu2Mg-MgZn2 alloys, shown in Fig. 8, have been measured

by [1964Kin], [1972Pre] and [1979Pre]. The Hf298 values of [1964Kin] are significantly more negative

than those of [1972Pre]. Furthermore, while the data of [1964Kin] show nearly smooth variation over the

whole composition range those of [1972Pre] exhibit a hump at 33.3 at.% Zn. [1972Pre] attributed this hump

to an interaction between the Fermi surface and Brillouin zones. On the other hand, in a later measurement

[1979Pre] obtained a concentration dependence of Hf298 that exhibits several minima and maxima, and it

is shown in Fig. 9. When the data of [1964Kin], [1972Pre]and [1979Pre] are compared, the latter two sets

show considerable scatter. However, [1988Pre] argued that the oscillating behavior of Hf298 should be

considered as a true alloying effect rather than a scatter even though both X-ray and neutron diffraction

[1979Ell] failed to identify any structural change or ordering within the homogeneity range of the C15

phase. However, if these oscillations really are true, in the Gibbs energy they must be compensated by the

term -T*S as G must be convex to lower values throughout the homogeneity range of the C15 Laves phase.

Since Zn atoms go into the Cu-sublattice, any peculiarity of the thermodynamic parameter of the C15 phase

is most likely to originate from the interaction between Cu and Zn atoms. A remarkable similarity between

the concentration dependence of Hf298 of the C15 phase and excess thermodynamic properties of the

Cu-Zn solid solution phases [1958Arg] was indeed noticed [1964Kin, 1988Pre].

[1987Hoc] calculated the enthalpy of formation of Cu2Mg, CuMgZn, CuMg2Zn3 and MgZn2 by model

considerations. The calculated [1987Hoc] and experimental [1964Kin, 1972Pre] values agree within

5 kJ mol-1. The low-temperature specific heat (both lattice and electronic components) of Cu2Mg-MgZn2

alloys was reported by [1967Ste] and [1971Bec].

A generalization of the Miedema model for the calculation of formation enthalpies of ternary and

higher-order intermetallics was developed by [1996Gon] and was successfully tested with respect to the

experimental data of alloys Mg(Cu1-xZnx)2.

Thermodynamic calculations (modelling) for the complete ternary diagram were presented by two groups

[1997Lia] and [1998Lia]. Mainly the two calculations agree fairly well, however the set of [1997Lia] is

much more simplified. These authors approximated Cu2Mg and MgZn2 as a single line compound with

constant Mg content of 33.33 at.% and as a stoichiometric compound. They neglected Cu-solubilities in

all binary Mg-Zn phases except MgZn2 and Mg-solubilities in , and of Cu-Zn.

References

[1906Gui] Guillet, M., “General Study of Special Bronze” (in French), Rev. Met., 3, 159-204 (1906)

(Experimental, 1)

[1934Lav] Laves, F., Löhberg, K., “The Crystal Structure of Intermetallic Compounds with Formula

AB2” (in German), Nachr. Ges. Wiss. Goettingen, 1, 59-66 (1934) (Crys. Structure,

Experimental,12)

[1936Lav] Laves, F., Witte, H., “The Influence of Valency Electrons on the Crystal Structures of

Ternary Magnesium Alloys” (in German), Metallwirtschaft, 15, 840-842 (1936) (Crys.


[1948Koe] Köster, W., Müller, F., “The Ternary System Copper-Zinc-Magnesium” (in German),

Z. Metallkd., 39, 352-359 (1948) (Equi. Diagram, Experimental, #, *, 6)

[1949Mik] Mikheeva, V.I., “On the Systematology of Liquidus Phase Diagrams of Ternary Metallic

Systems” (in Russian), Izv. Sekt. Fiz.-Khim. Anal., Akad. Nauk SSSR, 19, 126-133 (1949)

(Equi. Diagram, 7)

255


MSIT®

Cu–Mg–Zn

[1950Mik] Mikheeva, V.I., Kriyukova, O.N., “Melting Diagram of the System

Copper-Magnesium-Zinc” (in Russian), Izv. Sekt. Fiz.-Khim. Anal., Akad. Nauk SSSR, (20),


[1952Lie] Lieser, K.H., Witte, H., “Investigation of the Ternary Systems Magnesium-Copper-Zinc,

Magnesium-Nickel-Zinc and Magnesium-Copper-Nickel” (in German), Z. Metallkd., 43,


[1954Kle] Klee, H., Witte, H., “Magnetic Susceptibility of Ternary Magnesium Alloys and Its

Significance from the Point of View of Electron Theory of Metals” (in German), Z. Phys.

Chem. (Leipzig), 202, 352-378 (1954) (Experimental, 37)

[1954Lie] Lieser, K.H., Witte, H., “The Solubility of Hydrogen in Alloys I. Measurement and

Investigation of the Systems MgCu2-MgZn2 and MgNi2-MgZn2” (in German), Z. Phys.

Chem. (Leipzig), 202, 321-351 (1954) (Equi. Diagram, Experimental, 24)

[1955Gla] Gladyshevsky, E.I., Kripyakevich, P. I., “X-Ray Investigation of the System Cu-Mg-Zn in

the Region of MgCu2 -MgZn2” (in Russian), Nauk. Zap. Lviv. Derzhav. Univ. im. I.Franka,

Khim. Zbirnik, 34, 64-71 (1955) (Equi. Diagram, Experimental, #, *, 6)

[1956Gla1] Gladyshevsky, E.I., Kripyakevich, P.I., “The Solubility of Zn in Cu2Mg and Cu2Cd” (in

Russian), Izv. Sekt. Fiz.-Khim. Anal., 27, 209-211 (1956) (Equi. Diagram, Experimental, 6)

[1956Gla2] Gladyshevsky, E.I., Cherkashin, E.E., “Solid Solutions on the Basis of Metallic

Compounds”, Russ. J. Inorg. Chem., 22, 288-295 (1956) (Equi. Diagram, Experimental, 4)

[1956Mik] Mikheeva, V.I., “The Chemical Nature of the Ternary Intermetalic Phases in the Mg-Cu-Zn

and Mg-Cu-Ni Systems”, Proc. Acad. Sci., USSR, Chem. Sct., 109, 475-476 (1956)

(Experimental, 8)

[1957Sie] Siegelin, W., Leiser, K.H., Witte, H., “The Solubility of Hydrogen in Alloys III.

Investigation of the Ternary Systems MgCu2-MgAl2, MgCu2-MgSi2, MgNi2-MgCu2, and

the Binary System Ag-Cd, Cu-Be, Ag-Mg, Cu-Mg, Ni-Al, Ni-Si, Co-Al, Fe-Al” (in

German), Z. Elektrochem., 61, 359-376 (1957) (Experimental, 29)

[1958Arg] Argent, B.B., Wakerman, D.W., “Thermodynamic Properties of Solid Solutions I.

Copper-Zinc Solid Solution”, Trans. Faraday Soc., 54, 799-806 (1956) (Experimental,

Thermodyn., 21)

[1958Moe] Moeller, A., Witte, H., “The Temperature Dependence of Magnetic Susceptibility of the

Alloy System MgCu2-MgZn2” (in German), Z. Phys. Chem. (Frankfurt), 18, 130-132


[1960Wat] Watanabe, H., “On the Zn-Corner of the Quaternary System Zn-Al-Mg-Cu” (in Japanese),

J. Jpn. Inst. Met., 24, 672-676 (1960) (Equi. Diagram, Experimental, #, *, 3)

[1964Kin] King, R.C., Kleppa, O.J., “A Thermochemical Study of Some Selected Laves Phases”, Acta

Metall., 12, 87-97 (1964) (Experimental, Thermodyn., 27)

[1967Ste] Steiner, D., “The Specific Heat of the Alloys of the System MgCu2-MgZn2” (in German),

Z. Naturforsch. A, 22, 1284-1286 (1967) (Experimental, Thermodyn., 10)

[1970Kom] Komura, Y., Mitarai, M., Nakatani, I., Iba, H., Shimizu, T., “Structural Changes in the Alloy

Systems of Mg-Zn-Cu and Mg-Zn-Ag Related to the Friauf-Laves Phases”, Acta

Crystallogr., Sect. B: Struct.Crystallogr. Crys. Chem., 26, 666-668 (1970) (Crys. Structure,

Experimental, 8)

[1971Bec] Beckman, C.A., Craig, R.S., “Electronic Specific Heats of the MgCu2-xZnx System”,

J. Chem. Phys., 54, 898-901 (1971) (Experimental, Thermodyn., 10)

[1971Kri] Kripyakevich, P.I., Melnik, E.V., “Laves Phases with the Nine-Layer Structure in the

Systems Mg-Li-Zn, Mg-Cu-Zn and Mg-Co-Ni” (in Russian), Dopov. Akad. Nauk Ukr. RSR,

A(11), 1046-1048 (1971) (Crys. Structure, Experimental, 7)

[1971Sha] Shannette, G.W., Smith, J.F., “Single Crystalline Elastic Constants of Cubic MgCu2-MgZn2

Alloys”, J. Appl. Phys., 42, 2799-2803 (1971) (Experimental, 17)

[1972Pre] Predel, B., Ruge, H., “Study of the Enthalpies of Formation in the Systems Mg-Cu-Zn,

Mg-Cu-Al and Mg-Cu-Sn as a Contribution to the Understanding of the Binding Conditions

256


MSIT®

Cu–Mg–Zn

of Laves Phases” (in German), Mater. Sci. Eng., 9, 141-151 (1972) (Experimental,

Thermodyn., 61)

[1972Yam] Yamada, M., Matuki, K., “A Study of the Zn-rich Corner of the Zn-Cu -Mg Diagram” (in

Japanese), J. Jpn. Inst. Met., 36, 278-285 (1972) (Equi. Diagram, Experimental, #, *, 16)

[1975Nak] Nakaue, A., Inpue, K., Ikeda, T., “Pressure Effect on Stacking of Intermetallic Compound

Mg(Cu1-xZnx)2”, Proc. 4th Int. Conf. on High Pressure, Kyoto, Japan, 1974, 335-338 (1975)


[1977Oom] Oomi, G., “Composition Dependence of the Lattice Spacings of the Pseudobinary System

Mg(Cu1-xZnx)2 under High Pressure”, Jpn. J. Appl. Phys., 16, 1247-1248 (1977) (Crys.


[1977Ray] Raynor, G.V., “Constitution of Ternary and Some More Complex Alloys of Magnesium”,

Int. Met. Rev., 22, 65-96 (1977) (Crys. Structure, Review, 93)

[1979Cha] Chang, Y.A., Newmann, J.P., Mikula, A., Goldberg, D., “Cu-Mg-Zn”, in “Phase Diagrams

and Thermodynamic Properties of Ternary Copper-Metals Systems”, The Inter. Copper

Research Asso. Inc., New York (1979) (Equi. Diagram, Review, #, *, 14)

[1979Ell] Ellner, M., Predel, B., “Neutron Diffraction Investigation of Ternary Laves-phases of the

MgCu2-Type” (in German), J. Solid State Chem., 30, 209-221 (1979) (Crys. Structure,

Experimental, 26)

[1979Pre] Predel, B., Bencker, H., Vogelbein, W., Ellner, M., “Formation Enthalpies of Ternary Laves

Phases of the MgCu2-Type in the System MgCu2-MgZn2” (in German), J. Solid State

Chem., 28, 245-257 (1979) (Experimental, Thermodyn., 53)

[1984Nay] Nayeb-Hashemi, A.A., Clark, J.B., “The Cu-Mg (Copper-Magnesium) System”, Bull. Alloy

Phase Diagram, 5, 36-43 (1984) (Equi. Diagram, Review, #, *, 43)

[1987Hoc] Hoch, M., “Application of the Hoch-Arpshofen Model to the Thermodynamics of the

Cu-Ni-Sn, Cu-Fe-Ni, Cu-Mg-Al, and Cu-Mg-Zn Systems”, Calphad, 11, 237-246 (1987)

(Theory, Thermodyn., 16)

[1988Pre] Predel, B., “On the Thermochemistry of Intermetallic Compounds”, Thermochim. Acta,

129, 29-48 (1988) (Review, Thermodyn., 33)

[1996Gon] Gonçalves, A.P., Almeida, M., “Extended Miedema Model: Predicting the Formation

Enthalphies of Intermetallic Phases with More than Two Elements”, Physica, B228,

289-294 (1996) (Thermodyn., Theory, 19)

[1997Lia] Liang, H., Chang, Y.A., “A Thermodynamic Description for the Ternary Cu-Mg-Zn

System”, Z. Metallkd., 88, 836-841 (1997) (Thermodyn., Equi. Diagram, 31)

[1998Ans] Ansara, I., “COST 507, Thermochemical Database for Light Metal Alloys”, Ansara, I.,

Dinsdale, A.T., Rand, M.H. (Eds.), European Communities, Luxembourg, Vol. 2, (Cu-Mg)

170-174; (Cu-Zn) 186-191; (Mg-Zn) 227-233 (1998) (Equi. Diagram, Thermodyn.,

Assessment, 0)

[1998Lia] Liang, P., Seifert, H.J., Lukas, H.L., Ghosh, G., Effenberg, G., Aldinger, F.,

“Thermodynamic Modelling of the Cu-Mg-Zn Ternary System”, Calphad, 22, 527-544

(1998) (Thermodyn., 44)

257


MSIT®

Cu–Mg–Zn


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Cu)

< 1084.87

cF4

Fm3m

Cu

a = 361.3 pure Cu at 20°C

[V-C]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.944

c = 521.076

pure Mg at 25°C

[V-C]

(Zn)

< 419.58

hP2

P63/mmc

Mg

a = 266.47

c = 494.69

pure Zn at 25°C

[V-C]

CuMg2

< 568

oF48

Fddd

CuMg2

a = 907.0

b = 528.4

c = 1825

[V-C]

Cu2Mg

< 797

cF24

Fd3m

Cu2Mg

a = 703.4 at 33.33 at.% Mg

, CuZn (h)

902 - 454

cI2

Im3m

W

a = 299.67 36.1 to 55.8 at.% Zn

[V-C]

', CuZn (r)

< 468

cP2

Pm3m

CsCl

a = 295.39 44.8 to 50 at.% Zn

[V-C]

, Cu5Zn8

< 834

cI52

I43m

Cu5Zn8

a = 887.8 59.8 to 70.6 at.% Zn

[V-C]

, Cu0.7Zn2

700 - 560

hP3

P6

Cu1-xZn2

a = 427.5

c = 259

72.45 to 76 at.% Zn

[V-C]

, CuZn4

< 598

hP2

P63/mmc

Mg

a = 274.18

c = 429.39

78 to 88 at.% Zn

[V-C]

Mg7Zn3 (Mg51Zn20)

325 - 342

oI158

Immm

Mg51Zn20

a = 1408.3

b = 1448.6

c = 1402.5

[V-C]

MgZn

< 347

- - -

Mg2Zn3 (Mg4Zn7)

< 416

mC110

C2/m

Mg4Zn7

a = 2596.0

b = 524.0

c = 2678.0

[V-C]

258


MSIT®

Cu–Mg–Zn


MgZn2

< 590

hP12

P63/mmc

MgZn2

a = 522.3

c = 856.6

a = 518.4

c = 846.2

[V-C]

at MgCu0.11Zn1.89

[1971Kri]

Mg2Zn11

< 381

cP39

Pm3

Mg2Zn11

a = 855.2 [V-C]

* , CuMg2Zn3

< 718

hP24

P63/mmc

MgNi2

a = 512.4

c = 1682.0

[V-C]


Cu Mg Zn

L Cu2Mg 859 max L, Cu2Mg 39.2 33.2 27.6

L (Cu) + Cu2Mg 727 e1, max L 71.6 19.4 9.0

L + Cu2Mg 718 p3, max L 9.6 33.2 57.2

L + (Cu) Cu2Mg + 702 U1 L 51.7 13.2 35.1

L + Cu2Mg + 685 U2 L 38.8 12.0 49.2

L + MgZn2 603 p5, max L 1.0 33.4 65.6

L + Cu2Mg + 553 U3 L 12.0 16.0 72.0

+ + L 547 E1 L 12.7 7.3 80.0

L + + 514 U4 L 12.0 9.9 78.1

L Cu2Mg + (Mg) 448 e6, max L 9.8 77.8 12.4

L Cu2Mg + CuMg2 +(Mg) 440 E2 L 14.0 77.8 8.2

L + Cu2Mg + (Mg) 405 U5 L 2.3 74.8 22.9

L + + Mg2Zn11 400 P1 L 2.7 8.5 88.8

L + MgZn2 + Mg2Zn11 399 U6 L 2.1 8.9 89.0

L + MgZn2 + Mg2Zn3 376 U7 L 0.2 68.4 31.4

L + Mg2Zn11 + (Zn) 372 U8 L 1.4 6.4 92.2

L + + Mg2Zn3 MgZn 371 P2 L 0.2 68.8 31.0

L + (Mg) + MgZn 349 U9 L 0.2 71.5 28.3

L (Mg) + Mg7Zn3 + MgZn 341 E3 L 0.01 71.0 29.0

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

259


MSIT®

Cu–Mg–Zn

[1979Ell]

Zn, at.%

0

700

705

710

715

720

10 20 30 40 50

La

ttic

eP

ara

me

ter,

pm

Cu Mg2

[1955Glal]

[1952Lie]

[1971Sha]

Fig. 1: Cu-Mg-Zn.

Effect of Zn on the


the C15 Laves phase

along Cu2Mg -

MgZn2 section

10 20 30 40 50 60

400

500

600

700

800

900

Mg 33.33

Cu 66.67

Zn 0.00

Mg 33.33

Cu 0.00

Zn 66.67Zn, at.%

Te

mp

era

ture

, °C

Cu2Mg

L

τMgZn2

L+Cu2Mg

L+τ

τ+MgZn2Cu2Mg+τ

Fig. 2: Cu-Mg-Zn.

Approximately

pseudobinary section

Cu2Mg - MgZn2,

calculated from the

dataset of [1998Lia]

260


MSIT®

Cu–Mg–Zn

10 20 30 40 50 60

400

500

600

700

800

900

Mg Mg 0.00

Cu 38.46

Zn 61.54Zn, at.%

Te

mp

era

ture

, °C

L

(Mg)

Cu2Mg

L+Cu2Mg+γ

Cu2Mg+γ

L+Cu2Mg

L+Cu2Mg+(Mg)

(Mg)+Cu2Mg+traceCuMg2

L+(Mg)

Fig. 3: Cu-Mg-Zn.

Approximately

pseudobinary section

Mg - Cu5Zn8,

calculated from the


261


MSIT®

Cu–Mg–Zn

Fig

. 4a:

Cu

-Mg

-Zn

.R

eact

ion s

chem

e, c

alcu

late

dfr

om

the

dat

aset

of

[1998L

ia]

Cu

-Mg

Cu

-Zn

Mg-Z

nC

u-M

g-Z

n

l C

u2M

g +

(C

u)

72

5e 2

l +

(C

u)

β9

02

p1

l +

βγ

83

5p

2

l +

γδ

69

9p

4

L +

Cu

2M

g

τ7

18

p3

L +

τ M

gZ

n2

60

3p

5

LC

u2M

g +

CuM

g2

44

8e 6

l C

u2M

g +

CuM

g2

55

2e 4

δγ

+ ε

55

9e 3

l C

uM

g2 +

(M

g)

48

7e 5

L(M

g)+

Cu

2M

g+

Cu

Mg

24

40

E2

δ L

+ γ

+ ε

54

7E

1

L +

Cu

2M

g

γ +

τ5

53

U3

L +

β C

u2M

g +

γ6

85

U2

L +

(C

u)

Cu

2M

g +

β7

02

U1

L +

γτ

+ ε

51

4U

4

l +

δε

60

0p

6

(Cu

)+β+

Cu

2M

gL

+β+

Cu

2M

g

L+

γ+C

u2M

gβ+

γ+C

u2M

g

L+

γ+τ

L+

γ+ε

Cu

2M

g+

γ+τ

γ+ε+

τ

(Mg

)+C

uM

g2+

Cu

2M

g

L+

ε+τ

P1

U5

U6

U7

U5

262


MSIT®

Cu–Mg–Zn

Fig

. 4b

: C

u-M

g-Z

n.

Rea

ctio

n s

chem

e, c

alcu

late

dfr

om

the

dat

aset

of

[1998L

ia]

Cu

-Mg

Cu

-Zn

Mg-Z

nC

u-M

g-Z

n

lM

gZ

n +

Mg

7Z

n3

34

1e 8

l +

(M

g)

Mg

7Z

n3

34

1p

11

L(M

g)+

MgZ

n+

Mg

7Z

n3

34

1E

3

L +

τ(M

g)

+ M

gZ

n3

49

U9

L+

τ+

Mg

2Z

n3

MgZ

n3

71

P2

L+

ε(Z

n)+

Mg

2Z

n1

13

72

U8

L+

MgZ

n2

τ+M

g2Z

n3

37

6U

7

L+

τM

g2Z

n1

1+

MgZ

n2

39

9U

6

l +

MgZ

n2

Mg

2Z

n1

1

38

1p

9

lM

g2Z

n1

1 +

(Z

n)

36

7e 7

l +

Mg

2Z

n3

MgZ

n

34

7p

10

Mg

7Z

n3

MgZ

n+

(Mg)

32

5e 9

l +

MgZ

n2

Mg

2Z

n3

41

6p

8

L +

Cu

2M

g

τ +

(Mg)

40

5U

5

L +

τ +

ε M

g2Z

n1

14

00

P1

l +

ε (

Zn)

42

1p

7

L+

ε+M

g2Z

n1

1

ε+τ+

Mg

2Z

n1

1

Cu

2M

g+

τ+(M

g)

L+

τ+M

g2Z

n1

1

τ+M

gZ

n2+

Mg

2Z

n1

1

L+

τ+M

gZ

n2

L+

τ+M

g2Z

n3

MgZ

n2+

τ+M

g2Z

n3

L+

τ+(M

g)

L+

τ+M

gZ

nτ+

Mg

2Z

n3+

MgZ

n

ε+(Z

n)+

Mg

2Z

n1

1

L+

(Mg

)+M

gZ

n

τ+(M

g)+

Mg

Zn

L+

ε+τ

L+

Cu

2M

g+

Cu

Mg

2

L+

τ+M

gZ

n2

e 6p

3

L+

Cu

2M

g+

τ

p5

p5

U4

263


MSIT®

Cu–Mg–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Cu

Zn Data / Grid: mass%

Axes: mass%

p1

β

p2

U2

γ

U3

p4δ

U4

p6ε

p7

(Zn)

U5

p3

τ

p5

Cu2Mg

E2

e4

(Mg)(Cu)

600

500

500

600

850

700950

1000800

U1

E1

Mg2Zn

11

MgZn2

CuMg2

e5

e2

850

750

550

800

650

e6

10

20

10 20

80

90

Mg 25.00

Cu 0.00

Zn 75.00

Mg 0.00

Cu 25.00

Zn 75.00


Axes: at.%

(Zn)

e7

p7

p9

Mg2Zn

11

MgZn2

τE

1

U8

U4

ε p6

δ

γ

p4

U6 P

1

600

550

500

450

700

650

600

550

500

450

Fig. 5: Cu-Mg-Zn.

Liquidus surface,

calculated from the


Fig. 6: Cu-Mg-Zn.

Zn corner of the

liquidus surface,

calculated from the


264


MSIT®

Cu–Mg–Zn

40

700

600

500

400

300

200

Zn, at.%Cu

7

6

5

4

3

2

1

0

0 10 20 30

U1

Mg,at.%

Cu Mg2

e2

β

p1

Cu

Mg

Zn

60.00

0.00

40.00

0.30

0

0.15

0.20

0.25

0.05

0.10

0 0.80.4 1.2 2.01.6

Mg Zn2 11

350

330

310

290

270

250

230210

p4

�

390

370

U8

e7

Cu, at.%Zn

Mg

,a

t.%

170

150

190

Cu

Mg

Zn

2.00

0.00

98.00

Fig. 7: Cu-Mg-Zn.

Solvus surface of the

(Cu) solid solution,

calculated from the


Fig. 8: Cu-Mg-Zn.

Solvus surface of the

(Zn) solid solution,

calculated from the


265


MSIT®

Cu–Mg–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Cu


Axes: at.%

(Cu)+Cu2Mg

β

γ

CuMg2

Cu2Mg+CuMg

2

Cu 2

Mg+

CuM

g 2+(M

g)Cu 2M

g+(Mg)

Cu2Mg

(Zn)

ε

(Cu)

Mg2Zn

7

MgZn2

Mg2Zn

3

MgZn

(Mg)

Cu2Mg+γ

ε+τ

τ

(Mg)+

τ

Cu 2M

g+β

Fig. 9: Cu-Mg-Zn.


340°C, calculated

from the dataset of

[1998Lia]

266


MSIT®

Cu–Ni–Ti

Copper – Nickel – Titanium

Julius C. Schuster and Gabriele Cacciamani

Literature Data

In most experimental investigations Cu-Ni-Ti alloys were prepared from high purity metals (typically

99.9 %, e.g. [1969Sin, 1978Loo, 1980Yak1, 1986Yak, 1989Vav]) by melting under inert atmosphere

followed by heat treatment at various temperatures for various durations (e.g. 72 h at 1200°C [1969Sin],

100 h at 900°C [1986Yak], 144 h at 870 or 800°C [1978Loo], 240 h at 750°C [1969Sin], 50-300 h at

1000-1200°C [1994Jia]). It should be mentioned, that some authors used very short annealing times (e.g.

2 h at 840°C [1968Pfe]), which casts doubt, if thermal equilibrium was reached. [1978Loo] formed joints

between alloys of different compositions within the ternary Cu-Ni-Ti system and heat treated these

diffusion couples at 800 or 870°C for up to 900 h. The amorphous alloys investigated by [1989Vav,

1993Ali1, 1993Ali2, 1997Pus1, 1997Pus2, 2000Lou] were generally prepared by melt spinning eventually

followed by appropriate heat treatments in order to study the crystallization process.

Experimental techniques adopted in the study of this system and the compositions ranges investigated by

each author are summarized in Table 1.

Binary Systems

The binary systems are accepted from the MSIT Binary Evaluation Program: Cu-Ni [2002Leb], Cu-Ti

[2002Ans], and Ni-Ti [2003Ted]. Crystal structure data of the intermediate phases are given in Table 2.

Solid Phases

A description of the various solid Cu-Ni-Ti phases is reported in Table 2.

TiNi-based phase (CsCl type) undergo, at temperatures slightly above ambient, martensitic

transformations which form the basis for the so called “shape memory effect” (SME). The involved phases

are (orthorhombic, AuCd- type) and (monoclinic TiNi type), usually denoted in literature as B19 and

B19' respectively. These structures have been investigated by several authors [1981Zak, 1982Erm,

1985Che, 1993Lo, 1996Fuk, 1997Pus1, 1997Pus2, 1997Pus3, 1997Pus4, 2000Lou, 2000Vor, 2001Pot] as

a function of temperature and composition. [2001Pot] concluded that the orthorhombic phase forms from

the cubic CsCl type structure by atomic shifts in the (010)ORT layers and it differs from the AuCd type due

to a shift of the atoms from the centro-symmetric positions. The transformation to the monoclinic structure

is due to a monoclinic distortion of each second of the previous layers and, according to [2000Vor], it occurs

by a gradual accumulation of the degree of monoclinicity with the decreasing temperature, rather than

sharply. The anorthic structure determined by [1982Ero] was not confirmed by the following investigations

of the same composition and temperature range.

The phase Ti2(Ni1-xCux)3 (h) (x = 0.0583) forms by solid state reaction below 850°C [1990Nis]. Thus

[1978Loo] observed this phase in alloys equilibrated and quenched from 800°C, but not in alloys quenched

from 870°C. At 120°C this phase undergoes a second order transformation into Ti2(Ni1-xCux)3 (r)

(x = 0.0583) [1990Nis].

The phase Ti(Ni1-xCux)2 with the composition TiNiCu melts congruently at 1190°C [1980Yak2]. It is

reported to be stable at all temperatures investigated [1980Yak2, 1978Loo, 1968Pfe, 1992Ali2]. At 800°C

and 870°C a homogeneity range over a wide Cu/Ni ratio was found by [1968Pfe, 1978Loo]. For the same

phase at room temperature [1980Yak2] reports a single phase field within a much narrower Cu/Ti ratio but

having a width of about 4 at.% in the ratio (CuNi)/Ti.

Upon substituting Ni by Cu in TiNi3 [1966Vuc] observed four ternary phases at 2.5 at.% Cu (Ti10Ni29Cu),

at 5 at.% Cu (Ti5Ni14Cu), and two modifications (Cu3Sb type upon rapid quenching; TiAl3 type as cast) at

25 at.% Cu (TiNi2Cu). Ti10Ni29Cu was confirmed in as cast alloys by [1969Sin], but was not observed in

alloys annealed and quenched from 870 or 800°C [1968Pfe, 1978Loo]. Ti5Ni14Cu was found to exist from

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Cu–Ni–Ti

700°C up to 1090°C [1968Pfe, 1969Sin, 1978Loo]. No information on the melting behavior is available,

though Ti5Ni14Cu was observed in as cast alloys [1969Sin]. TiNi2Cu (Cu3Sb type) was corroborated in an

alloy Ti25Ni43Cu32 annealed and quenched from 930°C [1968Pfe], but was not seen in alloys equilibrated

at and quenched from 870 or 800°C [1968Pfe, 1978Loo]. TiNi2Cu (TiAl3 type) was corroborated in alloys

as cast [1968Pfe, 1969Sin], as well as equilibrated at 870°C [1978Loo], but not in alloys annealed and

quenched from 800°C [1968Pfe, 1978Loo].

Extended exchange of Cu by Ni or vice versa are reported by [1978Loo] for all binary phases at 870°C and

800°C (see Figs. 4 and 5 for details), by [1964Ram] for Ti2Cu (up to 25 at.% Ni), and by [1984Che1,

1984Che2] for Ti2Cu (5.2-7.7 at.%) and Ti2Ni (4.3-9.8 at.% Ni). The solution of Ni in TiCu2 stabilizes this

phase to lower temperatures [1968Pfe, 1978Loo].


Several vertical sections have been presented in literature as pseudobinary: Cu-TiNi3, Cu-TiNiCu,

TiNiCu-TiNi, TiNiCu-TiNi3, TiCu-TiNiCu, Ti2Cu-Ti2Ni, etc. However most of them cannot be of this type

without contradiction with some other phase equilibrium data. Only for the TiCu-TiNi section good

agreement among several authors can be found in literature; thus only this one is reported here (Fig. 1) as

assessed by [2000Tan2] even if, according to this assessment, what is generally described as a pseudobinary

eutectic is actually a very narrow three-phase field.

It is possible that some of the remaining sections (Cu-TiNi3 and Cu-TiNiCu-TiNi in particular) are not far

from a pseudobinary behavior (i.e. tie lines cross these sections with small angles).

Stable and metastable equilibria in the TiNi-TiCu section have been investigated by several authors

[1986Yak, 1986Ali, 1993Ali2, 1997Pus3, 1997Pus4, 2000Tan2], especially in the Ni-rich region, due to the

interest for the SMA. A Calphad assessment of the section, including the low temperature and phases,

has been carried out by [2000Tan1, 2000Tan2].

In this section a eutectic equilibrium occurs at 960°C. Most authors agree with this temperature. The value

of 924°C determined by [1986Yak] from DTA cooling curves seems unreliable, also considering the

disagreement with the accepted melting temperatures of the TiCu and TiNi phases. The solubility of in

TiCu is not exceeding 3 at.% Ni [1978Loo, 1986Ali, 1986Yak]. The solubility limit of TiCu in is much

higher and quite uncertain: 32 at.% Cu and 27 at.% Cu at the eutectic temperature were reported by

[1986Yak] and [1986Ali], respectively; [1978Loo] measured 33 at.% Cu at 870°C and 800°C. According

to the optimization by [2000Tan1] the solubility of TiCu in increases quite sharply with temperature when

approximating the eutectic temperature (this could explain the poor agreement among the experimental

determinations).

In this section, moreover, amorphous and nano-crystalline phases may be formed by fast cooling (melt

spinning technique). A summary of the metastable behavior of these alloys is given in [1997Pus3]. Finally,

the martensitic transitions taking place in the Ni-rich side of the section and responsible for the

thermoplastic properties of the SMA have been studied by different authors [1979Bri, 1979Mer, 1981Zak,

1982Erm, 1982Ero, 1985Che, 1986Edm, 1987Plo, 1987Tok, 1990Tsu, 1991Tsu, 1993Lo, 1993Mat2,

1997Pus1, 1997Pus2, 1997Pus3, 1997Pus4, 2000Lou, 2000Vor, 2001Pot]. A summary of their results is

here reported in Fig. 2 according to [1997Pus4].


The invariant equilibria here accepted are listed in Table 3 and a partial reaction scheme is presented in

Fig. 3.

The partial system TiCu-TiNi-Ti is probably characterized by one ternary eutectic (E3, 860°C) and two

U type reactions (U1 and U3) involving the liquid phase. These have been here proposed adapting quite

contradictory literature information [1968Tak, 1979Bud, 1981Bud2, 1982Yak, 1983Kov, 1984Che2,

1992Ali2] to the accepted binaries and mass balance rules. The invariant equilibrium related to the

( Ti) ( Ti) transition was either reported as a U type at 780°C [1981Bud2] or an E type at 738°C

[1982Yak]: the second one is here selected because the authors present more detailed results.

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Cu–Ni–Ti

In the partial system Cu-TiCu-TiNi the equilibria proposed by [1971Fed, 1980Yak1, 1992Ali3] are

generally accepted with minor variations in order to be consistent with the accepted 870°C isothermal

section by [1978Loo].

In the partial system Cu-Ni-TiNi a more complex system of invariant equilibria should be present, related

to the different ternary phases which eventually form in the solid state. A ternary eutectic is probably present

at >70 at.% Ni and about 1025°C [1971Fed] but it is not clear which solid phases take part in it. Closer to

the Cu-Ni binary system a eutectic valley connects the L Ni+TiNi3 binary eutectic to e2 [1967Jos2].

Liquidus Surface

A complete set of monovariant liquid curves was reported by [1992Ali3], and sets of liquidus isothermal

curves by [1968Tak, 1982Yak]. Even in this case contradictory results are present in literature. The

incomplete liquidus surface shown in Fig. 4 is the result of the combination of the data found in [1967Jos2,

1968Tak, 1971Fed, 1979Bud, 1980Yak1, 1980Yak2, 1981Bud2, 1982Yak, 1983Kov, 1986Ali, 1986Yak,

1988Ali2, 1992Ali1, 1992Ali2, 1992Ali3] with the already accepted binary and ternary equilibria.

Isothermal Sections

Isothermal sections over the entire composition range were investigated by [1968Pfe] at 800°C (reproduced

by [1968Pfe], [1979Dri] and [1987Jen]) and by [1978Loo] (reproduced in [1987Jen]) at 800 and 870°C.

The results of [1978Loo] (Figs. 5 and 6) are preferred, because [1968Pfe] used only 2 hours annealing time

for alloy equilibration, while [1978Loo] used equilibration times of 1 week or more. The partial isothermal

sections at 1000 and 850°C by [1971Fed] (reproduced in [1971Fed]) show in detail the homogeneity limit

of Cu coexisting with TiNi3 and supposedly TiNiCu (TiNi is given instead). Neither CuNi14Ti5 nor

CuNi2Ti are indicated. The isothermal section at 600°C presented by [1992Ali2] was apparently derived

from data originated by as cast alloys and thus will not be considered here.

According to the tie lines shown in these sections neither Ni-TiNiCu nor TiCu-TiNiCu can be pseudobinary

sections, as claimed by [1992Ali1] and [1986Ali], respectively.

The partial isothermal sections of the Ti-rich corner for temperatures between 1015 and 700°C [1980Sha,

1981Bud2] are theoretical constructions based on the assumption of a pseudobinary system Ti2Cu-TiNi2 as

well as on a Cu-Ti binary diagram different from that accepted here. Thus these isotherms are not considered

further.


Some of the temperature-composition sections not accepted as pseudobinary are presented here. In fact they

are probably quite reliable at least at higher temperatures, as far as the liquid is involved in the equilibria.

In Figs. 7 and 8 the Cu-TiNi3 and Cu-TiNi sections are presented according to [1967Jos2] and [1980Yak2,

1986Ali, 1986Yak], respectively, with minor corrections in order to make them coherent with the accepted

equilibria.

Thermodynamics

The enthalpies of transformation between the different low symmetry phases in the TiNi-TiCu section up

to 25 at.% Cu were determined by DSC measurements by [2000Tan2] and are reported in Fig. 9. The

transformation enthalpy of Cu-Ni-Ti shape memory alloys was also measured by [1992Yu].


The importance of the Cu-Ni-Ti alloys is related to the SME (shape memory effect) quite recently found in

the TiNi based alloys with Cu as main addition element. Materials showing this effect are commonly

denoted as SMA (shape memory alloys) and they are still under study and continuous improvement. See

[1997Fuk, 1999Ots, 1999Roe, 1999Sil, 2000Wu] for a recent review on the SMA properties and their

present and potential applications.

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Cu–Ni–Ti

Several mechanical properties of selected Cu-Ni-Ti alloys were studied: shear strength [1968Tak], tensile

strength [1972Dea, 1987Har], fatigue [1972Dea], hardness [1972Dea, 1973Kni, 1984Che1, 1984Che2].

Miscellaneous

The effect of neutron radiation on the mechanical properties of an alloy at Cu-5.3Ni-3.3Ti (at.%) was

determined by [1987Har]. [1984Wil] refer that addition of Ti reduces the capability of the Cu-Ti alloys to

dissolve oxygen. Alloying with 3-8 mass% Cu improves the wetting of diamond by Ni-Ti alloys [1984Che1,

1984Che2]. A Ti-18Ni-12Cu (at.%) alloy resulted to be resistant to corrosion by HNO3 or H2SO4, but

mildly attacked by HCl [1968Tak]. Phase diagram studies in the quaternary system Cu-Ni-Si-Ti were

carried out by [1981Bud1, 1984Ali, 1985Ali1, 1985Ali2, 1988Ali1, 1988Ali2, 1990Ali, 1991Ali,

1992Ali4].

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270


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Cu–Ni–Ti

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Mater. Trans., JIM, 37(10), 1540-1546 (1996) (Experimental, 15)

[1997Fuk] Fokuda, T., Saburi, T., Kakeshita, T., Kitayama, M., “Shape Memory Behaviour of

Ti-40.5Ni-10Cu Alloy Affected by C11b-Type Precipitates”, Mat. Trans, JIM, 38, 107-111


[1997Pus1] Pushin, V.G., Volkova, S.B., Matveeva, N.M., “Structural and Phase Transformations in

Quasi-binary TiNi-TiCu Alloys Rapidly Quenched from the Melt: I. High-Copper

Amorphous Alloys”, The Phys. Metals Metallogr., 83, 275-282 (1997) (Experimental, 14)

[1997Pus2] Pushin, V.G., Volkova, S.B., Matveeva, N.M., “Structural and Phase Transformation in

Quasibinary TiNi-TiCu Alloys Rapidly Quenched from the Melt: II. Alloys with Mixed”,

The Phys. Metals Metallogr., 83, 283-288 (1997) (Experimental, 8)

273


MSIT®

Cu–Ni–Ti

[1997Pus3] Pushin, V.G., Volkova, S.B., Matveeva, N.M., “Structural and Phase Transformations in

Quasi-Binary TiNi-TiCu Alloys Rapidly Quenched from the Melt: III. Mechanisms of

Crystallization”, The Phys. Metals Metallogr., 83, 435-443 (1997) (Experimental, 7)

[1997Pus4] Pushin, V.G., Volkova, S.B., Matveeva, N.M., Yurchenko, L.I., Chistyakov, A.S.,

“Structural and Phase Transformations in Quasi-Binary TiNi-TiCu Alloys Rapidly

Quenched from Melt: VI. Martensitic Transformations”, Phys. Met. Metallogr., 84, 441-448


[1999Ots] Otsuka, K., Ren, X., “Recent Developments in the Research of Shape Memory Alloys”,

Intermetallics, 7, 511-528 (1999) (Review, 131)

[1999Roe] Roesner, H., Shelyakov, A.V., Glezer, A.M., Feit, K., Schlossmacher, P., “Study of an

Amorphous-Crystalline Structured Ti-25Ni-25Cu (at.%) Shape Memory Alloys”, Mater.

Sci. Eng., A273-275, 733-737 (1999) (Experimental, 11)

[1999Sil] Silvain, J.F., Chazelas, J., Lahaye, M., Trombert, S., “The Use of Shape Memory Alloy NiTi

Particles in SnPbAg Matrix: Interfacial Chemical Analysis and Mechanical

Characterisation”, Mater. Sci. Eng., A273-275, 818-823 (1999) (Experimental, 15)

[2000Lou] Louzguine, D.V., Inoue, A., “Crystallization Behavior of Ti50Ni25Cu25 Amorphous Alloy”,

J. Mater. Sci., 35, 4159-4164 (2000) (Experimental, 8)

[2000Tan1] Tang, W., Sandstroem, R., Miyazaki, S., “Phase Equilibria in the Pseudobinary

Ti0.5Ni0.5-Ti0.5Cu0.5 System”, J. Phase Equilib., 21(3), 227-234 (2000) (Thermodyn.,

Calculation, 37)

[2000Tan2] Tang, W., Sandstrom, R., Wei, Z.G., Miyazaki, S., Verhoeven, J.D., “Experimental

Investigation and Thermodynamic Calculation of the Ti-Ni-Cu Shape Memory Alloys”,

Metall. Mat. Trans., 31A, 2423-2430 (2000) (Experimental, Thermodyn., 22)

[2000Vor] Voronin, V.I., Naish, V.E., Novoselova, T.V., Pushin, V.G., Sagaradze, I.V., “Structures of

Monoclinic Phases in Titanium Nickelide: I. Transformation Cascade B2-B19-B19”, Phys.

Met. Metallogr., 89(1), 12-18 (2000) (Crys. Structure, Experimental, 10)

[2000Wu] Wu, S.K., Lin, H.C., “Recent Development of TiNi-Based Shape Memory Alloys in

Taiwan”, Mater. Chem. Phys., 64, 81-92 (2000) (Review, 102)

[2001Pot] Potapov, P.L., Shelyakov, A.V., Schryvers, D., “On the Crystal Structure of TiNi-Cu

Martensite”, Scr. Mater., 44, 1-7 (2001) (Crys. Structure, Experimental, 15)

[2002Ans] Ansara, I., Ivanchenko, V., “Cu-Ti (Copper-Titanium)”, MSIT Binary Evaluation Program,

in MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services GmbH,

Stuttgart; Document ID: 20.11457.1.20, (2002) (Equi. Diagram, Review, 26)

[2002Leb] Lebrun, N., “Cu-Ni (Copper - Nickel)”, MSIT Binary Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH,

Stuttgart; Document ID: 20.14832.1.20, (2002) (Crys. Structure, Equi. Diagram,

Assessment, 51)

[2003Ted] Tedenac, J-C., Velikanova, T., Turchanin, M., “Ni-Ti (Nickel-Titanium)”, MSIT Binary


International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi.

Diagram, Assessment, 37)

Table 1: Recent Investigations of the Cu-Ni-Ti System

Reference Experimental Techniques Temperature/Composition/Phase Range

[1966Vuc] XRD TiNi3

[1967Jos1] XRD, LOM, SEM+EPMA TiNi3

[1967Jos2] XRD, LOM, SEM+EPMA Cu-TiNi3 section

[1968Pfe] XRD Isothermal section at 800°C

274


MSIT®

Cu–Ni–Ti

[1968Tak] TA, Hardness Partial liquidus surface

[1969Sin] XRD, LOM Ti(Ni1-xCux)3

[1971Fed] XRD, LOM, DTA Cu-rich corner at 700-1000°C

up to 3 at.% Ti and/or Ni

[1973Kni] TEM Tix(Ni3Cu7)1-x 0 x 0.02

[1978Loo] Diffusion couples, XRD, LOM,

SEM+EPMA

Isothermal sections at 800 and 870°C

[1978Mel] TEM Ti50Ni50-xCux 0 < x < 35

[1979Bri] XRD, TEM Ti50Ni50-xCux 0 < x < 25

[1979Bud] XRD, DTA, LOM, Hardness, Resistivity Ti-Ti2Ni-Ti2Cu at 700-1000°C

[1979Mer] Electrical resistivity Ti50Ni50-xCux 0 < x < 30

[1980Yak1] XRD Cu-rich corner

up to 28 at.% Ti and 44 at.% Ni

[1980Yak2] XRD, LOM, SEM+EPMA Cu-rich corner

up to 28 at.% Ti and 44 at.% Ni

[1981Zak] XRD, LOM, Deformation Ti50Ni50-xCux 10 < x < 40

[1982Erm] XRD, LOM, Deformation Ti50Ni50-xCux 10 < x < 33

[1982Ero] XRD, Resistivity, Deformation Ti50Ni40Cu10

[1982Yak] XRD, LOM, SEM+EPMA, DTA Ti-rich corner

up to 42 at.% Cu and 45 at.% Ni

[1986Ali] XRD, LOM,DTA TiCu-TiNi-TiNiCu at 800°C

[1986Edm] Electrical resistivity Ti50Ni45Cu5

[1986Yak] LOM, SEM+EPMA, DTA TiCu-TiNi section

[1987Plo] Acoustic emission Ti50Ni50-xCux 0 < x < 11

[1989Vav] XRD, DTA, Resistivity Amorphous alloys, various comp.

[1990Nis] XRD, LOM, TEM Ti40Ni60-xCux 0 < x < 50

[1990Tsu] DSC Ti50Ni50-xCux 6 < x < 9

[1991Tsu] Deformation Ti50Ni50-xCux 6 < x < 9

[1992Ali1] XRD, LOM, DTA Liquidus surface in the Cu-Ni-TiNi

triangle

[1992Ali2] XRD, LOM, DTA Liquidus surface, isothermal section at

600°C

[1992Ali4] XRD, LOM, DTA Liquidus surface in the TiNi-TiCu-Ti2Ni

triangle

[1992Gou] TEM, DTA, DSC, IF Ti55Ni42Cu3

[1993Ali1] XRD, SEM, DTA Stable and metastable equilibria in the

Ti-rich corner,

0-20 at.% Cu, 0-20 at.% Ni

[1993Ali2] XRD, SEM+EPMA, DTA,

Hardness

TiNi-TiCu section

amorphous alloys


275


MSIT®

Cu–Ni–Ti


[1993Lo] DSC, XRD, IF,

Electrical resistivity

Ti50Ni40Cu10

-100 < T < 100°C

[1993Mat1] Mechanical properties Ti50Ni25Cu25

[1994Jia] Diffusion couples, SEM+EPMA Cu-Ni-Ti at 1000-1200°C

0-25 at.% Ti, a few at.% Cu

[1997Pus1] XRD, Electron diffraction, TEM,

Electrical resistivity, Stress/strain,

Microhardness

Ti50Ni50-xCux 0 < x < 40

[1997Pus2] XRD, Electron diffraction, TEM Ti50Ni50-xCux 10 < x < 20



Ti50Ni50-xCux 0 < x < 40



Ti50Ni50-xCux 0 < x < 40

[2000Lou] XRD, Electron diffraction,

TEM, DSC

Ti50Ni25Cu25

[2000Tan1] CALPHAD TiNi-TiCu vertical section

[2000Tan2] DSC,

CALPHAD

Ti50Ni50-xCux 0 < x < 13

0 < x < 25

[2000Vor] Neutron diffraction Ti50Ni40Cu10 and Ti50Ni25Cu25

[2001Pot] XRD+Rietveld

TEM

Ti50Ni25Cu25

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

, (Cu1-x Nix)1-y Tiy

Cu

< 1085

Ni

< 1455

cF4

Fm3m

Cu a = 352.4

a = 356.50

a = 361.46

0 x 1, 0 y < 0.14

x = 0, y = 0

x = 0.52, y = 0

x = 1, y = 0

[Mas2, V-C2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 dissolves 14 at.% Cu

and 10 at.% Ni [Mas2]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

dissolves 1.6 at.% Cu

and 0.2 at.% Ni [Mas2]

TiCu4

< 885

oP20

Pnma

Au4Zr

a = 453.0

b = 434.2

c = 1293.0

78-80.9 at.% Cu

dissolves 1 at.% Ni

[Mas2, V-C2]


276


MSIT®

Cu–Ni–Ti

Ti(Cu1-xNix)2

Cu2Ti

890 - 870

oC12

Amm2

Au2V

a = 442

b = 800

c = 451

a = 436.3

b = 797.7

c = 447.8

x = 0.167, T = 870°C,

stable at T = 800°C

[V-C2]

x = 0

Ti2Cu3

< 875

tP10

P4/nmm

Cu3Ti2

a = 313

c = 1395

dissolves 5 at.% Ti

[V-C2]

Ti3Cu4

< 925

tI14

I4/mmm

Ti3Cu4

a = 312.6

c = 1996.4

dissolves 7 at.% Ti

[V-C2]

TiCu

< 982

tP4

P4/nmm

TiCu

a = 310.7 to 312.5

c = 588.7 to 591.9

48-52 at.% Cu

dissolves 2 at.% Ti

[Mas2, V-C2]

Ti2(Cu1-xNix)

Ti2Cu

< 1005

tI6

I4/mmm

MoSi2

a = 294

c = 1067

a = 294.38

c = 1078.6

0 x 0.4

x = 0.25 [V-C2]

x = 0

TiNi3< 1380

hP16

P63/mmc

TiNi3

a = 510.88

c = 831.87

dissolves 5 at.% Cu

[V-C2]

, Ti(Ni1-xCux) (h)

TiNi

1310 - ~50

cP2

Pm3m

CsCl

a = 302.6

a = 298.6

a = 301.2

0 x 0.4

x = 0.2 [1982Ero]

43-50.5 at.% Ti at x = 0

45 at.% Ti, x = 0

at 51 at.% Ti, x = 0 [V-C2]

´, Ti(Ni1-xCux) (l1)

70

oP4

Pmma

AuCd

(also denoted

B19)

a = 299

b = 398

c = 450

a = 290.2

b = 429.0

c = 451.8

a = 291.77

b = 429.00

c = 450.39

x = 0.2, described as AuCd type

[1982Ero]

x = 0.5, described as AuCd type

[2000Vor]

x = 0.5, described as AuCd derivative

[2001Pot]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

277


MSIT®

Cu–Ni–Ti

´´, Ti(Ni1-xCux) (r)

70

mP4

P21/m

NiTi

(AuCd derivative,

also denoted

B19 )

a = 288.9

b = 412

c = 462.2

= 96.8°

a = 290

b = 426

c = 457

= 97.5°

a = 288.6

b = 422.0

c = 452.8

= 90.80°

x = 0 [1971Ots]

x = 0.18 [1979Bri]

x = 0.2

= 90.80° at T = 300 K

= 94.81° at T = 250 K

= 96.08° at T = 150 K

= 96.23° at T = 78 K

[2000Vor]

Ti2(Ni1-xCux)

< 984

cF96

Fd3m

NiTi2

a = 1131.93 x = 0 - 0.15

[V-C2]

, Ti(Ni1-xCux) h** also denoted R phase,

related to phase.

Forms for x(Ni+Cu) > xTi

[1985Che, 1992Gou]

* 1, Ti(Ni1-xCux)2

TiNiCu

< 1190

tI6

I4/mmm

MoSi2

a = 310

c = 798

a = 314

c = 800

x = 0.20 - 0.82 [1978Loo]

x = 0.23 [V-C2]

x = 0.75 [V-C2]

* 2, Ti2(Ni1-xCux)3 (h)

850 - 120

tP10 a = 440.28

c = 1352.5

related to Ti2Cu3

x = 0.0583 [1978Loo]

* 3, Ti2(Ni1-xCux)3 (r)

< 120

o** a = 462.0

b = 452.2

c = 1319.4

x = 0.0583 [1990Nis]

* 4, Ti5Ni14Cu

1090 - 700

hR12

R3m

BaPb3

a = 511.2

c = 1887.1

[1966Vuc];

cited in [V-C2] as CuNi13Ti5

* 5, TiNi2Cu oP8

Pmmn

Cu3Sb

a = 506.7

b = 421.6

c = 449.5

possibly a high temperature phase

[V-C2]

* 6, TiNi2Cu tI8

I4/mmm

Al3Ti(Ni3V)

a = 361.1

c = 745.9

a = 311.4

c = 795.5

possibly a high temperature phase

from a sample at Cu50Ni25Ti25

[1966Vuc]

* Ti10Ni29Cu hP40

P63/mmc

Pd2RhTa

a = 511.6

c = 2090

possibly a metastable phase [V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

278


MSIT®

Cu–Ni–Ti



Ti Ni Cu

L TiNiCu 1190 congruent L

TiNiCu

33.3

33.3

33.3

33.3

33.4

33.4

L TiNiCu + TiNi 1140 e2 (max) L

TiNiCu

TiNi

40

35.5

49.4

40

35.5

49.4

20.0

30.0

1.2

L TiNiCu + TiNi3 1100 e4 (max) - - - -

L (Cu) + TiNiCu 1070 e5 (max) L

(Cu)

TiNiCu

6.5

-

-

6.5

3.7

31.4

87

-

-

L (Cu) + TiNiCu + TiNi3 1025 E1 - - - -

L TiNiCu + TiNi3 (?) + TiNi 1000 E2 L 35 46 19

L TiNiCu + TiCu 980 e7 (max) L 45.5 9 45.5

L TiCu + TiNi 960 e8 (max) L 50 48 2

L TiNiCu + TiCu + TiNi 930 E3 L 48 10 42

L + TiCu Ti2Cu + TiNi 930 U1 L 60 10 30

L + TiCu TiNiCu + Ti3Cu4 913 U2 L 37.5 3.5 59

L + ( Ti) Ti2Cu + Ti2Ni 900 U3 L 66 21 13

L + (Cu) TiNiCu + TiCu4 > 870

< 885

U4 L 23 3 74

L + Ti3Cu4 TiNiCu + Cu2Ti 873 U5 L 31 3 66

L TiNiCu + TiCu4 + TiCu2 > 870

< 873

E4 L 26.5 3 70.5

L Ti2Cu + TiNi + Ti2Ni 860 E5 L 65.5 22.5 12

( Ti) Ti2Cu + Ti2Ni + ( Ti) 738 E6 ( Ti)

( Ti)

93

100

5

0

2

0

279


MSIT®

Cu–Ni–Ti

10 20 30 40

-250

0

250

500

750

1000

1250

Ti 50.00

Ni 50.00

Cu 0.00

Ti 50.00

Ni 0.00

Cu 50.00Cu, at.%

Te

mp

era

ture

, °C

α

960

L

TiCu

α´ + TiCu

α´´ + TiCuα´´

α´

1310

10 20 30

0

Ti 50.00

Ni 50.00

Cu 0.00

Ti 50.00

Ni 10.00

Cu 40.00Cu, at.%

Te

mp

era

ture

, °C

Af´

As´

TR

Ms´

Mf´ α +

α''

α' +

α''

α''

α'

Mf

As

Ms

Af

α

Fig. 2: Cu-Ni-Ti.

Martensitic

transformations in the

alloys of the

pseudobinary

TiNi-TiCu section

obtained by rapid

solidification at

105 K s-1

[1997Pus4]

Fig. 1: Cu-Ni-Ti.

The pseudobinary

TiNi-TiCu section

including the low

temperature and

phases [2002Tan2]

280


MSIT®

Cu–Ni–Ti

Fig. 3: Cu-Ni-Ti. Partial reaction scheme

Cu-Ti A-B-C

L (Cu)+TiNiCu

1070 e5max

L TiNi+TiNiCu

1140 e2max

Cu-Ni-Ti Ni-Ti

L TiCu+TiNiCu

980 e7max

L TiNi+TiNiCu+TiNi3

1000 E2

L TiCu+TiNi+TiNiCu930 E3

L+γ TiCu4+TiNiCu>870<885 U

4

L+Ti3Cu

4 TiCu

2+TiNiCu873 U

5

L TiCu2+TiCu

3+TiNiCu>870<873 E

4

l+TiCu Ti3Cu

4

925 p2

l + (Cu) TiCu4

885 p4

l+Ti3Cu

4TiCu

2

890 p3

l TiCu2+TiCu

4

875 e11

l (Ni) + TiNi3

1304 e1

l TiNi + TiNi3

1118 e3

TiCu+TiNi+TiNiCu

?

L γ+TiNiCu+TiNi3

1025 E1

L TiNiCu+TiNi3

1100 e4max

?

L+TiCu Ti2Cu+TiNi~930 U

1

L+(βTi) Ti2Cu+Ti

2Ni~900 U

3

L TiNi+Ti2Ni+Ti

2Cu~860 E

5

(βTi) (αTi)+Ti2Cu+Ti

2Ni738 E

6

l TiCu+Ti2Cu

960 e9

l (βTi)+Ti2Cu

1005 e6

(βTi) (αTi)+Ti2Cu

790 e12

l+TiNi Ti2Ni

948 p1

l Ti2Ni + (βTi)

942 e10

(βTi) (αTi)+Ti2Ni

765 e13

TiNi+Ti2Ni+Ti

2Cu

Ti2Cu+Ti

2Ni+(αTi)

L+TiCu Ti3Cu

4+TiNiCu913 U

2

L TiCu+TiNi

960 e8max

TiCu+TiNi+Ti2Cu

281


MSIT®

Cu–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni

Cu Data / Grid: at.%

Axes: at.%

(βTi)(αTi)

Ti2Ni TiNi TiNi

3

Ti2Cu

TiCu

Ti3Cu

4

Ti2Cu

3

TiCu2

βTiCu4

τ1

TiNi2Cu

τ4

γ

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni


Axes: at.%

Ti2Ni

E5

U3

α

(βTi)

U1

p1

Ti2Cu

TiCue

8max

E3

e7max

E1e

2max

TiNi3

γ

τ1

U2

Ti3Cu

4

U5

TiCu2

E4

U4

TiCu4

Cu

e5max

p4

e11

p2

e9

p3

e6

e1

e10

e4max

e3

E2

Fig. 5: Cu-Ni-Ti.


870°C [1978Loo]

Fig. 4: Cu-Ni-Ti.

Partial liquidus

surface

282


MSIT®

Cu–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni


Axes: at.%

γ

TiNi3TiNiTi

2Ni

τ2

τ4

τ1

(αTi)(βTi)

Ti2Cu

TiCu

Ti3Cu

4

Ti2Cu

3

TiCu4

TiNixCu

2-x

20 40 60

1000

1100

1200

1300

1400

Cu Ti 25.00

Ni 75.00

Cu 0.00Ni, at.%

Te

mp

era

ture

, °C

1060

1085°C

1380°C

TiNi3γ

L

Fig. 6: Cu-Ni-Ti.


800°C [1978Loo]

Fig. 7: Cu-Ni-Ti.

Polythermal section

Cu-Ni3Ti [1967Jos2]

283


MSIT®

Cu–Ni–Ti

10 20 30 40

1000

1100

1200

1300

1400

Cu Ti 50.00

Ni 50.00

Cu 0.00Ni, at.%

Te

mp

era

ture

, °C

1070

1190

1085°C

1140

1310°C

ατ1

γ

-2000

-1800

-1600

-1400

-1200

-1000

-800

-600

-400

-200

0

0 5 10 15 20 25

Cu, at.%

Enth

alp

y,(J

mol

)DH

·-1

��´´

��´

��´´

Fig. 8: Cu-Ni-Ti.

Polythermal section

Cu-NiTi [1980Yak2,

1986Ali]

Fig. 9: Cu-Ni-Ti.

Calculated enthalpy H

for the

transformations

´´, ´ and

´ ´´

[2000Tan2]

284


MSIT®

Cu–Si–Ti

Copper – Silicon – Titanium

Natalia Bochvar, Yong Du, Dmitriy Kevorkov, Rainer Nast and Peter Rogl

Literature Data

Information on the phase relations within the Cu-Si-Ti ternary essentially are due to the investigations of

several groups of authors supplying (a) data on the isothermal section at 800°C [1967Nic, 1969Nic,

1974Spr] (b) on the solidification behavior within a predominantly pseudobinary section Ti5Si3-Ti2Cu

[1977Bud] and in the two adjacent regions: Ti-Ti5Si3-Ti2Cu [1975Bud, 1976Ali, 1978Bud, 1982Bud] and

TiCu-Ti2Cu-Ti5Si3 [1985Ali] as well as (c) on the join TiCu-TiSi [1992Lut], and (d) within the polythermal

sections for the Ti-rich part at 0.2, 1, 2 and 3 mass% Cu [1976Nar]. Crystallographic studies concerned with

the ternary compounds are due to [1967Nic, 1969Nic, 1970Nic, 1974Spr]. For their investigations

[1967Nic, 1969Nik, 1970Nic, 1974Spr] employed powder and single crystal X-ray diffraction techniques,

LOM and EMPA on as cast as well as annealed alloy specimens (24 to 100 h at 800 to 1100°C). Samples

were argon arc-melted from the elements with a purity better then 99.6% via the formation of binary Cu-Ti

master alloys. Ternary compounds were also synthesized from Cu/Si-rich melts removing the Cu/Si-phases

afterwards in diluted HNO3 [1967Nic, 1969Nic, 1970Nic, 1974Spr]. Alloys in the region around the

Ti5Si3-Ti2Cu pseudobinary were arc- or levitation-melted and homogenized by subsequent annealing for

up to 500 h in order to reach equilibrium [1975Bud, 1976Ali, 1977Bud, 1985Ali]; [1992Lut] used arc

melted alloys which were annealed at 800°C for 200 h. Techniques of investigation comprised LOM,

EMPA, XRD and microhardness measurements. Reviews on the Cu-Si-Ti system are due to [1979Cha] and

[1979Dri].

Binary Systems

The binary boundary system Cu-Si and Cu-Ti were accepted from [2002Leb] and [2002Ans]. For Si-Ti

critical assessment and thermodynamic modelling is due to [1996Sei]. Crystallographic data for the

boundary phases are compiled in Table 1. From the experimental description of the Cu-Si-Ti system it

seems that the binary phase Ti2Cu shows a non-negligible homogeneity region in contrast to the line

compound in the binary version in [2002Ans].

Solid Phases

Two ternary phases were established from the isothermal section at 800°C [1969Nic], TiCuSi with

practically no ternary range of homogeneity and TiCu0.1Si1.9, which may essentially be considered as a

stabilization of the ZrSi2 type structure by small amounts of copper. In addition to these phases a further

compound TiCuSi2 was mentioned from titanium silicide crystal growth experiments in copper flux

(1300°C, quench) [1972Jan]. TiCuSi2 was suggested to be stable at higher temperatures but was only

obtained in small quantities from hot-pressed samples [1972Jan]. Crystallographic data of the ternary

compounds are listed in Table 1.

Solid solubility of Si in copper-titanium compounds at 800°C was reported to be generally small, not

exceeding 1-2 at.% Si [1974Spr]. However, from the investigation of the Ti5Si3-Ti2Cu pseudobinary the

authors of [1975Bud, 1977Bud, 1978Bud, 1982Bud] arrived at a solubility of ~2.7 at.% Si in Ti2Cu at

800°C. At higher temperatures the solubility of Ti5Si3 in Ti2Cu rises rapidly up to 26 at.% Si at 1600°C, the

temperature of the pseudobinary peritectic formation of Ti2Cu [1975Bud]. The rather large solubility of

4 at.% Si and 6 at.% Ti in (Cu) as shown in the 800°C section by [1979Cha] was probably taken from a

preliminary study of [1970Nic], whereas a later and more detailed investigation by [1974Spr] arrived at 6.5

at.% Si and 1.8 at.% Ti. These values are in slight contrast to the accepted maximum binary solubility of

about 11 at.% Si and of about 4 at.% Ti in (Cu) at 800°C. The values shown in the ternary isothermal section

at 800°C in Fig. 1 were changed in order to comply with the accepted binaries. Except for Ti5Si3, which

according to [1975Bud] dissolves up to 1.9 at.% Cu, solid solubilities of Cu in titanium silicides were shown

285


MSIT®

Cu–Si–Ti

to be very small [1974Spr] i.e. 0.7 at.% Cu in Ti3Si [1975Bud]. The solubility of Cu+Si in ( Ti) was

reported not to exceed 1.5 mass% [1975Bud]. Solid solubilities of Ti in copper silicides at 800°C were not

specified in detail but were shown in the isothermal section by [1974Spr] (see Fig. 1).


The vertical section Ti5Si3-Ti2Cu investigated by [1975Bud] is shown in Fig. 2. This section reveals a

pseudobinary peritectic reaction (maximum): L+Ti5Si3 Ti2Cu(Si) at 1600°C (see Table 2 and Fig. 2). For

some inconsistencies in the low temperature solubility of Si in Ti2Cu, as reported by [1974Spr] (< 1 at.%

Si) and [1975Bud] (2.7 at.% Si), see section Solid Phases.


The isothermal reactions reported (see Table 2) concern a pseudobinary peritectic and some consecutive

solidification reactions descending into the adjacent sides of the Ti2Cu-Ti5Si3 pseudobinary [1975Bud] and

[1976Ali, 1985Ali]. The partial reaction scheme is shown in Fig. 3.

Liquidus Surface

A partial liquidus surface in the Ti-rich corner of the ternary (see Fig. 4) was established from the

investigations by the authors of [1975Bud] and [1985Ali].

Isothermal Sections

Figure 1 presents the complete isothermal section at 800°C revealing the existence of two ternary

compounds. The location of TiCuSi2, observed by [1972Jan] as a likely high-temperature phase, not

encountered at 800°C [1974Spr], is shown as a small half filled circle. The authors of [1976Ali, 1977Bud]

presented a series of nine partial isothermal sections covering the range Ti-Ti2Cu-Ti5Si3 at 810, 830, 870,

900, 950, 1010, 1100, 1200 and 1300°C: the evolution of phase relations with temperature is shown in Figs.

5 to 13 with small adaptions to comply with the accepted phase diagrams. The main discrepancy arises from

the fact that the binary phase Ti2Cu seems to have a small homogeneity region in the Cu-Ti binary and

further in the ternary. As a consequence the vertices of the three phase region Ti+Ti2Cu+Ti5Si3 are shown

by [1976Ali, 19787Bud] at considerably lower Si content than expected from the pseudobinary section

Ti2Cu-Ti5Si3 of [1975Bud]. Figures 5 to 13 follow the values given by [1976Ali, 1978Bud] to be taken at

the Ti-poor phase boundary of Ti2Cu at a concentration slightly smaller than 33.3 at.% Cu. Similarly a set

of partial isothermal sections was constructed by [1985Ali] for the region TiCu-Ti2Cu-Ti5Si3 to portray the

phase relations in the temperature region from 915 to 1900°C. The section at 935°C (Fig. 14) follows the

maximum solid solubility of Si in Ti2Cu (at the Ti-poor phase boundary at 66.67 at.% Ti) corresponding to

the pseudobinary system of [1975Bud].

Thermodynamics

Activities of Si in binary Cu-Si and ternary Cu-Si-Ti melts were measured at 1550°C using the isopiestic

technique where the activity coefficient of Si Si was directly determined from the reaction

(SiO2)+2H2 2H2O(g)+(SiCu) setting up an equilibrium among the slags, Cu-Si-Ti melt and a gas phase

controlling the oxygen potential in the gas phase via a gas mixture H2O/H2 of constant composition. The

analytical expression for the activity coefficient of Si in binary Cu-Si melt at 1550°C was

ln Si = -5.69+6.69xSi-26.22xSi2 (see Fig. 15), for an oxygen potential (PH2O/PH2)2 = 1.34 10-3 the mole

fraction of Si in the melt at 1550°C xSiTi varies almost linearly with the mole fraction of Ti as

xSiTi 102 = 5.87+185xTi (see Fig. 16) [1994Xue1, 1994Xue2].

From the experimental data, the activity in reaction coefficients of Si in the melts have been determined and

activity and reaction coefficients of Ti in binary Cu-Ti melts at 1550°C have been estimated [1994Xue1,

1994Xue2].

286


MSIT®

Cu–Si–Ti


A series of isopleths was constructed for the Ti-rich corner at 0.2, 1, 2 and at 3 mass% Cu based on

differential thermal and microstructural analyses, electrical resistivity, Vickers hardness and high

temperature strengths [1976Nar]. These isopleths, however, are with respect to sequence and temperature

regions of the various phases observed inconsistent with the isothermal sections independently investigated

by [1974Spr] and [1976Ali, 1977Bud] and thus are not presented here.


A Vickers hardness of 700 kp mm-2 was reported for the (001) face of TiCuSi [1967Nic].

Miscellaneous

Phase relations were extended to the system TiCu-Ti5Si3-TiCuNi by [1988Ali] and to TiCu-TiNi-TiCuNi

by [1990Ali].

References

[1967Nic] Nickl, J.J., Sprenger, H., “Single Crystals of the Ternary Compounds TiCuSi and ZrCuSi”

(in German), Naturwissenschaften, 54, 18 (1967) (Crys. Structure, Experimental, 1)

[1969Nic] Nickl, J.J., Sprenger, H., “About New Compounds in the Titanium-Copper-Silicon Ternary

System” (in German), Z. Metallkd., 60, 136-139 (1969) (Crys. Structure, Equi. Diagram,

Experimental, 11)

[1970Nic] Nickl, J.J., Schweitzer, K.K., “Gasphase Metallurgical Investigation of the Ti-Si System”

(in German), Z. Metallkd., 61, 54-61 (1970) (Crys. Structure, Experimental, 31)

[1972Jan] Jangg, G., Kieffer, R., Blaha, A., Sultan, T., “Solubility of Disilicides in Auxiliary Metal

Melts” (in German), Z. Metallkd., 63, 670-676 (1972) (Experimental, 24)

[1974Spr] Sprenger, H., “The Ternary Systems (Titanum, Zirconium, Hafnium)-Copper-Silicon” (in

German), J. Less-Common Met., 34, 39-71 (1974) (Crys. Structure, Equi. Diagram,


[1975Bud] Budberg, P.B., Alisova, S.P., “Phase Diagram of the Ti-Ti5Si3-Ti2Cu System” (in Russian),

Dokl. Akad. Nauk SSR, 224(1), 157-159 (1975) (Equi. Diagram, Experimental, #, *, 3)

[1976Ali] Alisova, S.P., Budberg, P.B., “Phase Transformations in the System Ti-Ti5Si3 - Ti2Cu”,

Russ. Metall. (Engl. Transl.), 5, 182-188 (1976) (Equi. Diagram, Experimental, 4)

[1976Nar] Nartova, T.T., Zuykova, N.A., “Phase Equilibria and Properties of Ti-Cu-Si Alloys”, Russ.

Metall. (Engl. Transl.), 2, 166-169 (1976) (Equi. Diagram, Experimental, *, 5)

[1977Bud] Budberg, P.B., Alisova, S.P., “Study of Alloys of the Ti-Cu-Si System” (in Russian),

Metalloved. Legkikh. Splavov, 265-271 (1977) (Equi. Diagram, Experimental,*, 7)

[1978Bud] Budberg, P.B., Alisova, S.P., “Phase Diagram of the Ti-Ti5Si3-Ti2Cu System” (in Russian),

Titan- Metalloved. Tekhnol. Tr. 3-1, Mezhdunar. Konf: Titan., Moskow (2), 409-412 (1978)

(Equi. Diagram, Experimental, *, 3)

[1979Cha] Chang, YA, Neumann, J.J., Mikula, A., Goldberg, D., “Cu-Si-Ti”, INCRA Monograph, Ser.

6, Phase Diagram and Thermodynamic Properties of Ternary Copper-Metals Systems”,


[1979Dri] Dritz, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L.,

Turkina, N.I., “Binary and Multicomponent Copper-Base Systems” (in Russian), Nauka,

Moskow, 153-156 (1979) (Review, 4)

[1982Bud] Budberg, P.B., Alisova, S.P., “Constitutional Diagram of the Ti-Ti5Si3-Ti2Cu System”,

Titanium Alloys, Proc. Int. Conf., 3rd, 1976, 2, 1351-1355 (1982) (Equi. Diagram,

Experimental, *, 79)

[1985Ali] Alisova, S.P., Budberg, P.B., Kobylkin, A.N., “Phase Transformation in Alloys of the

System Ti2Cu-TiCu-Ti5Si8”, Russ. Metall., 1, 201-203 (1985), translated from Izv. Akad.

Nauk SSR, Met., 1, 197-199 (1985) (Equi. Diagram, Experimental, #, 6)

287


MSIT®

Cu–Si–Ti

[1988Ali] Alisova, S.P., Kovneristyi, Yu.K., Budberg, P.B., “Phase Structure of Titanium Copper

(TiCu)-Titanium Silicon (Ti5Si3)-Titanium Copper Nickel (TiCuNi(T)) Melts and

Characteristic of the Supercooled State” (in Russian), Rasplavy, 2(2), 104-106 (1988) (Equi.


[1990Ali] Alisova, S.P., Budberg, P.B., Kovneristyi, Yu.K., “Five-Phase Transformation in the

Four-Fold System of Titanium Intermetallic Compounds with the Participation of Liquid

Phase” (in Russian), Rasplavy, 4, 83-87 (1990) (Equi. Diagram, Experimental, 6)

[1992Lut] Lutskaya N.V., Alisova S.P., “Phase Structure of Ti-Cu(Ni,Co)-TiSi Sections in Ternary

Ti-Cu(Ni,Co)-Si Systems”, Russ. Metall., 3, 180-182 (1992), translated from Izv. Akad.

Nauk SSR, Met., 3, 194-196 (1992) (Equi. Diagram, Experimental, #, 9)

[1994Ole] Olesinski, R. W., Abbaschian, G. J., Cu-Si (Copper-Silicon), in “Phase Diagrams of Binary

Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM

International, Materials Park, OH, 398-405 (1994) (Review, Equi. Diagram, Crys.

Structure, Thermodyn., #, *, 60)

[1994Sub] Subramanian, P.R., “Cu (Copper)”, in “Phase Diagrams of Binary Copper Alloys”,

Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E., (Eds.), ASM International, Materials

Park, OH, 1-3 (1994) (Crys. Structure, Thermodyn., Review, 16)

[1994Xue1] Xue, X., Che, Y., “Activity Interaction Coefficients of Si in Cu-Ti-Si Melts at 1550°C”,

Z. Metallkd,. 85(6), 391-393 (1994) (Thermodyn., Experimental, 17)

[1994Xue2] Xue, X., Che, Y., Du, H., “Activity Interaction Coefficients of Si in Cu-Ti-Si Melts at

1550°C”, J. Mater. Sci. Technol., 10, 67-70 (1994) (Thermodyn., Experimental, 17)


Z. Metallkd., 87, 2-13 (1996) (Review, Thermodyn., Equi. Diagram, *, 65)

[2002Ans] Ansara, I., Ivanchenko, V., “Cu-Ti (Copper-Titanium)”, MSIT Binary Evaluation Program,

in MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services GmbH,

Stuttgart; Document ID: 20.11457.1.20, (2002) (Equi. Diagram, Review, 26)

[2002Leb] Lebrun, N., Dobatkina, T., Kuznetsov, V., Li, C., “Cu-Si (Copper-Silicon)”, MSIT Binary





Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Cu)

< 1084.62

cF4

Fm3m

Cu

a = 361.46 at 25°C [Mas2]

melting point [1994Sub]

( Si) hP4

P63/mmc

La

a = 380

c = 628

at 25°C, 16 GPa 1 atm [Mas2]

( Si) cI16

Ia3

Si

a = 663.6 at 25°C, 16 GPa [Mas2]

( Si) tI4

I41/amd

Sn

a = 468.6

c = 258.5

at 25°C, 9.5 GPa [Mas2]

288


MSIT®

Cu–Si–Ti

( Si)

< 1414

cF8

Fd3m

C (diamond)

a = 543.06 at 25°C [Mas2]

( Ti) hP3

P6/mmm

Ti

a = 462.5

c = 281.3

at 25°C, HP 1 atm [Mas2]

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 [Mas2]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35

at 25°C [Mas2]

, Cu7Si

842 - 552

hP2

P63/mmc

Mg

a = 256.05

c = 418.46

11.05 to 14.5 at.% Si [Mas2, V-C2]

, Cu6Si

852 - 785

cI2

Im3m

W

a = 285.4

14.2 to 17.2 at.% Si [Mas2]

at 14.9 at.% Si [1994Ole]

, Cu5Si(h)

824(?) - 710(?)

t** a = 881.5

c = 790.3

17.6 to 19.6 at.% Si [Mas2],

sample was annealed at 700°C (?) and

analyzed at 850°C (?) [V-C2]

, Cu5Si(r)

< 729

cP20

P4132

Mn

a = 619.8 17.15 to 17.6 at.% Si [Mas2, 1994Ole]

Cu15Si4800

cI76

I43d

Cu15Si4

a = 961.5 21.2 to 21.3 at.% Si [Mas2, V-C2]

, Cu3Si(h2)

859 - 558

hR*

Rm

or

t**

a = 247

= 109.74°

a = 726.7

c = 789.2

23.3 to 24.9 at.% Si [Mas2, 1994Ole]

[V-C2]

’, Cu3Si(h1)

620 - 467

hR*

R

a = 472

= 95.72°

23.2 to 25.2 at.% Si [Mas2, 1994Ole]

’’, Cu3Si(r)

< 570

o** a = 7676

b = 700

c = 2194

23.3 to 24.9 at.% Si [Mas2, 1994Ole]

, Ti2Cu 1)

< 1012

tI6

I4/mmm

MoSi2

a = 295.3

c = 1073.4

[Mas2, V-C2, 1994Ole]

TiCu

< 982

tP4

P4/nmm

TiCu

a = 310.8 to 311.8

c = 588.7 to 592.1

48 to 52 at.% Cu [Mas2, V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

289


MSIT®

Cu–Si–Ti

Ti3Cu4

< 925

tI14

I4/mmm

Ti3Cu4

a = 313.0

c = 1994

[Mas2, V-C2]

Ti2Cu3

< 875

tP10

P4/nmm

Ti2Cu3

a = 313

c = 1395

[Mas2, V-C2]

TiCu2

890 - 870

oC12

Amm2

VAu2

a = 436.3

b = 797.7

c = 447.8

[Mas2, V-C2]

TiCu4

885 - 400

oP20

Pnma

ZrAu4

a = 452.5

b = 434.1

c = 1295.3

~78 to ~80.9 at.% Cu [Mas2, V-C2]

TiCu4

< 500

tI10

I4/m

MoNi4

a = 228.59

c = 358.45

~78 to ~80.9 at.% Cu [Mas2]

Ti3Si

< 1170

tP32

P42/n

Ti3P

a = 1026.6

c = 506.9

dissolves 0.7 at.% Cu at 800°C

[1975Bud, V-C2, Mas2]

Ti5Si3< 2130

hP16

P63/mcm

Mn5Si3

a = 746.1

c = 515.08

35.5 to 39.5 at.% Si, congruent point,

dissolves 1.9 at.% Cu at 800°C

[1975Bud, V-C2, Mas2]

Ti5Si4(h)

1920 - 1300

oP36

Pnma

Sm5Ge4

a = 650.6

b = 1269.0

c = 664.5

[1970Nic, Mas2]

Ti5Si4(l)

< 1300

tP36

P41212

Zr5Si4

a = 671.3

c = 1217.1

[1970Nic, Mas2]

TiSi (h)

1570 - 900

oP8

Pnma

FeB

a = 655.1

b = 363.3

c = 498.3

[V-C2, Mas2]

TiSi (l)

< 900

oC32

Cmmm

TiSi

a = 1874

b = 708.1

c = 359.6

[1970Nic]

TiSi2< 1500

oF24

Fddd

TiSi2

a = 826.71

b = 480.00

c = 855.05

congruent point

[V-C2, Mas2]

* 1, TiCuSi oP12

Pnma

Co2Si

a = 619.3

b = 374.6

c = 713.0

observed at 800°C

[1969Nic, V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

290


MSIT®

Cu–Si–Ti

1) Ti2Cu is likely to exhibit a small homogeneity region (see also sections “Solid Phases” and “Isothermal

Sections”).


*All compositions were read from the original experimental diagrams and were slightly modified to comply with

the accepted binary data.

* 2, Ti1(CuxSi1-x)2 oC12

Cmcm

ZrSi2

a = 356.2

b = 1353.1

c = 355.0

observed at 800°C

[1969Nic, V-C2]

* 3, TiCuSi2 ? ? formed from flux growth experiments

[1972Jan]

Reaction T [°C] Type Phase Composition*, at.%

Ti Cu Si

l + Ti5Si3 Ti2Cu ~1600 p1 (max) l

Ti5Si3Ti2Cu

65.4

63

63.5

22.4

1.5

10.5

12.2

35.5

26

L + Ti5Si3 ( Ti) + Ti2Cu 1200 U1 L

Ti5Si3( Ti)

Ti2Cu

72.5

64.3

95.1

64.7

21.8

2

1.5

18.4

5.7

33.7

3.4

16.9

( Ti) + Ti5Si3 Ti3Si + Ti2Cu 950 U2 ( Ti)

Ti5Si3Ti3Si

Ti2Cu

97.1

64.5

75.0

66.6

1.1

1.9

0.7

30

1.8

33.6

24.3

3.4

L + Ti2Cu Ti5Si3 + TiCu 935 U3 L 52 47 1

( Ti) + Ti3Si ( Ti) + Ti2Cu 830 U4 ( Ti)

Ti3Si

( Ti)

Ti2Cu

95

75.0

98.6

66.7

4.6

0.7

0.6

30.5

0.4

24.3

0.8

2.8

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

291


MSIT®

Cu–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Cu


Axes: at.%

TiSi2

Cu3Si + TiSi

2 + Si

τ1 + τ

2 + Cu3Si

τ1

τ1 + Ti

5Si3 + (Cu)

Ti5Si

3 + (Cu) + βTiCu4

Ti5Si

3 + Ti2Cu

3 + βTiCu4

TiCuTi2Cu(βTi)

(αTi)

Ti3Si

Ti5Si

3

Ti5Si

4(l)

TiSi(l) τ2

Ti5 Si

3 + TiCu + Ti

2 Cu(αTi) + Ti

3Si + Ti

2Cu

Ti3Cu

4Ti

2Cu

3βTiCu

4 (Cu)

Cu7Si

βδ

Cu15

Si4

η' (Cu3Si(h

1))

τ3

Fig. 1: Cu-Si-Ti.


800°C

10 20 30

1000

1250

1500

1750

2000

Ti 62.49

Cu 0.00

Si 37.51

Ti 66.70

Cu 33.30

Si 0.00Cu, at.%

Te

mp

era

ture

, °C

2130°C

1600

1012°C

Ti5Si3+ L

Ti2Cu + L

Ti2CuTi5Si3 + Ti2Cu

Ti5Si3

Fig. 2: Cu-Si-Ti.

Vertical section

Ti5Si3 - Ti2Cu

292


MSIT®

Cu–Si–Ti

Fig. 3: Cu-Si-Ti. Reaction scheme

Si-Ti A-B-CCu-Si-Ti Cu-Ti

l (βTi)+Ti5Si

3

1340 e1

(βTi)+Ti5Si

3Ti

3Si

1170 p1

(βTi) (αTi)+Ti3Si

860 e4

l+Ti5Si

3 Ti

2Cu

1600 p1(max)

L+Ti5Si

3(βTi)+Ti

2Cu1200 U

1

L+Ti5Si

3+Ti

2Cu

l (βTi)+Ti2Cu

1005 e2

(βTi)+Ti5Si

3Ti

3Si+Ti

2Cu950 U

2

Ti5Si

3+Ti

3Si+Ti

2Cu

(βTi)+Ti3Si+Ti

2Cu

L+Ti2Cu Ti

5Si

3+TiCu935 U

3

l Ti2Cu+TiCu

960 e3

(βTi)+Ti3Si (αTi)+Ti

2Cu830 U

4

L+Ti5Si

3+TiCu

Ti5Si

3+Ti

2Cu+TiCu

(αTi)+(βTi)+Ti2Cu

(βTi) (αTi)+Ti2Cu

790 e5

(βTi)+Ti2Cu+Ti

5Si

3

(αTi)+Ti3Si+Ti

2Cu

(βTi)+Ti2Cu+L

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%

Ti5Si

3(ζ)

(βTi)

Ti3Si(γ)

Ti2Cu(ε)

U1

1200

°C

950°C

830°C

Fig. 4: Cu-Si-Ti.

Projection of

invariant planes and

monovariant lines in

the partial

Ti-Ti5Si3-Ti2Cu

system, after

[1975Bud]

293


MSIT®

Cu–Si–Ti

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%Ti

5Si

3(ζ)

(βTi)

Ti2Cu(ε)

γ

(αTi)

γ+ζ+ε

(αTi) + (βTi) + γ + ε

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%

Ti5Si

3(ζ)

(βTi)

Ti2Cu(ε)

γ

(αTi)

γ+ζ+ε

(αTi)+γ+ε

(αTi)+(βTi)+ε

Fig. 6: Cu-Si-Ti.


830°C; after

[1976Ali]

Fig. 5: Cu-Si-Ti.


810°C; after

[1976Ali]

294


MSIT®

Cu–Si–Ti

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%

Ti5Si

3(ζ)

Ti2Cu(ε)

γ

(αTi)

γ+ζ+ε

(βTi)+γ+ε

(βTi)

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%Ti

5Si

3(ζ)

γ

(βTi)

Ti2Cu(ε)

(βTi)+γ+ε

γ + ζ + ε

Fig. 7: Cu-Si-Ti.


870°C; after

[1976Ali]

Fig. 8: Cu-Si-Ti.


900°C; after

[1976Ali]

295


MSIT®

Cu–Si–Ti

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%Ti

5Si

3(ζ)

Ti2Cu(ε)

(βTi)

γ

(βTi)+ζ+γ+ε

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%Ti

5Si

3(ζ)

γ

(βTi)

L

Ti2Cu(ε)

(βTi)+ε+L

(βTi)+ζ+ε

(βTi)+γ+ζ

L+ε

Fig. 9: Cu-Si-Ti.


950°C; after

[1976Ali]

Fig. 10: Cu-Si-Ti.

Isothermal section

system at 1010°C;

after [1976Ali]

296


MSIT®

Cu–Si–Ti

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%

Ti5Si

3(ζ)

Ti2Cu(ε)

L

L + ε

(βTi)

L +( βTi)+ε

(βTi)+ζ+γ

(βTi)+ζ+ε

γ

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%

(βTi)L

Ti5Si

3(ζ)

Ti2Cu(ε)

L+ε

L+(βTi)+ζ+ε

(βTi)+L

Fig. 11: Cu-Si-Ti.


1100°C; after

[1976Ali]

Fig. 12: Cu-Si-Ti.


1200°C; after

[1976Ali]

297


MSIT®

Cu–Si–Ti

70

80

90

10 20 30

10

20

30

Ti Ti 60.00

Cu 40.00

Si 0.00

Ti 60.00

Cu 0.00


Axes: at.%

Ti5Si

3(ζ)

(βTi)

L

L+(βTi)+ζ

L+ε

Ti2Cu(ε)

L+ζ

L+ζ+ε

50

60

70

80

90

10 20 30 40 50

10

20

30

40

50

Ti Ti 40.00

Cu 60.00

Si 0.00

Ti 40.00

Cu 0.00


Axes: at.%

L+ζ+ε+TiCu

L+ε+TiCu

L+ε+TiCuL

Ti5Si

3(ζ)

Ti2Cu TiCu

Fig. 13: Cu-Si-Ti.


1300°C; after

[1976Ali]

Fig. 14: Cu-Si-Ti.

Partial isothermal

section at 935°C; after

[1985Ali]

298


MSIT®

Cu–Si–Ti

ln = -5.69 + 6.69 -26.22(gSi x xSi Si2)

x Si × 102

0.20 0.40 0.60 0.80 1.00

-5.10

-5.30

-5.50

-5.70

lng S

i

0 1.20

xTi

Si Ti⋅ ⋅= 5.81 + 1.853x10 102 2r = 0.93

xTi ⋅ 102

0.50 1.50 2.50 3.00 3.502.001.00

xT

i

Si

⋅10

2

5.00

10.00

0.00

Fig. 15: Cu-Si-Ti.

Activity coefficient of

Si in binary Cu-Si

melt (ln Si) vs mole

fraction of Si (xSi) at

1550°C

Fig. 16: Cu-Si-Ti.

Mole fractions of Si

(xSiTi) in the Cu-Si-Ti

melt at 1550°C vs

mole fractions of Ti

(xTi).

r is the correlation

coefficient of the

regression equation

299


MSIT®

Fe–Ni–Ti

Iron – Nickel – Titanium

Gautam Ghosh

Literature Data

A fairly large number of experimental studies have been carried out to establish the ternary phase equilibria[1938Vog, 1941Vog, 1963Spe, 1967Dud, 1981Loo, 1999Abr]. The first comprehensive study of phaseequilibria was carried out by [1938Vog]. They used metallography, thermal analysis and X-ray diffractiontechniques. [1963Spe] reported partial isothermal sections, representing phase equilibria of Fe corner, at700 and 1100°C. [1967Dud] established a pseudobinary section, TiFe-TiNi, using twenty-one ternary alloysprepared using elements of purity greater than 99.94%. They used metallography, thermal analysis, X-raydiffraction techniques to determine the phase equilibria. [1981Loo] employed diffusion couples techniqueto determine the isothermal section at 900°C. They prepared alloys using elements of following purity:99.95% Fe, 99.99% Ni and 99.97% Ti. Three types of diffusion couples, element/element, element/alloy,alloy/alloy, and fifteen in total, were prepared by solid state resistance welding. The couples were thensealed in evacuated silica tubes and annealed at 900°C for up to 900 h. Except [1981Loo], otherexperimental studies were restricted to alloys containing less than 50 at.% Ti. The results of these phaseequilibrium studies were reviewed earlier [1985Gup, 1991Gup]. More recently, [1994Ali1] reported the phase equilibria of alloys containing more than 50 at.% Ti. Theyused Armco Fe, N-00 grade electrolytic Ni, and iodide Ti to prepare alloys by arc melting in an inertatmosphere. They prepared a number of ternary alloys in the Ni and Fe atomic ratios of 1:3, 1:1, and 3:1.The phase equilibria were determined using thermal analysis, metallography and X-ray diffraction.[1994Ali2] carried out rapid solidification of ternary alloys containing up to 33.85 at.% Fe and 26.8 at.%Ni. The alloys were subsequently annealed at 900°C for 25 h, and the microstructures were compared.[1994Jia] measured the partitioning of Fe and Ti between (Ni) and TiNi3 (or the tie-lines) at 1000, 1100 and1200°C. They prepared diffusion couples between 5Fe-Ni (mass%) and 21Ti-1Fe-Ni (mass%) alloys. Thecouples were annealed for up to 300 h followed by quenching into iced brine. The compositions of thephases were measured by electron probe microanalysis. [1999Abr] determined the isothermal section at 1000C using both bulk alloys and diffusion couples. Theyprepared bulk alloys using iodide grade Ti, electrolytic grade Ni and carbonyl grade Fe in an arc furnace.The bulk alloys were equilibrated at 1200°C for 150 h. They also prepared a large number of couples usingFe-Ni, Fe-Ti and Ni-Ti alloys that were welded at 1200°C and at a pressure of 19.6 MPa. The couples weresubsequently annealed at 1200°C. The diffusion paths and phase compositions were established by meansof electron probe microanalysis. Besides graphical representation, [1999Abr] tabulated the tie-line andtie-triangle compositions. Based on diffusion couple results, [1999Abr] identified six two-phase regions(TiNi-TiFe2, TiNi-TiNi3, TiNi3-TiFe2, (Fe,Ni)-TiNi3, (Fe,Ni)-TiFe2) and three three-phase regions( (Fe,Ni)-TiFe2-TiNi3, TiFe2-TiNi-TiNi3, ( Ti)-TiFe-Ti2Ni). [1999Efi] studied microstructures ofTiFe/Ni diffusion couples annealed at 1200°C for up to 1.5 h. Based on the dynamics of interfacialmicrostructure, they derived the interdiffusion coefficients. Besides phase equilibria studies at hightemperatures, the effect of Fe on the low temperature martensitic transformations of TiNi has also beenstudied extensively [1982Hwa1, 1982Hwa2, 1982Nis, 1985Goo, 1986Edm, 1992Shi, 1993Mat, 1995Gue,1997Har, 2000Chu, 2000Vor, 2000Xu]. Some of the results of recent phase equilibria studies werereviewed by [2001Gup].

Binary Systems

The Fe-Ti, Fe-Ni and Ni-Ti binary phase diagrams are accepted from [1991Mur], [2004Kuz] and[2004Ted], respectively.

300


MSIT®

Fe–Ni–Ti

Solid Phases

[1968Abr] reported that lattice parameters of two sets of austenitic alloys: Fe-27Ni (at.%) containing up to10 at.% Ti and Fe-30Ni (at.%) containing up to 6 at.% Ti. Powders of these alloys were solution treated at1024°C and quenched to room temperature. Variation of lattice parameter gave the solubility limit of Ti inaustenite at 1024°C. For example, it was about 5.4 at.% Ti in Fe-27Ni (at.%) alloy.The solubility of Fe in ( Ti) can be increased from 24 to 29.5 at.% by rapid solidification [1994Ali2] wherethe cooling rate was estimated to be 106 °C s-1. Rapid solidification also suppresses the eutectoid reaction( Ti) ( Ti)+TiFe. In the ternary alloys, up to 25.95 at.% Fe and 8.23 at.% Ni can be dissolved in ( Ti) byrapid solidification. TiFe and TiNi form a continuous solution in the solid state [1967Dud, 1981Loo], and the lattice parameterdecreases linearly from TiNi to TiFe [1967Dud]. At 900°C, TiFe2 dissolves up to 28 at.% Ni, TiNi3dissolves up to 14 at.% Fe, and Ti2Ni dissolves up to 26 at.% Fe [1981Loo]. The solubility of Fe in TiNi3increases to 15.4 0.8 at.% Fe at 1200°C [1999Efi]. In TiNi3 and NiTi2, Fe resides primarily on the Nisublattice. In TiFe2, Ni resides primarily on the Fe sublattice.Binary B2-TiNi undergoes martensitic transformation at low temperature to a monoclinic structure,commonly known as B19’ martensite [1992Shi, 1995Gue]. However, at slightly Ni-rich composition, thepresence of dislocations, or the presence of metastable Ti3Ni4 precipitates are known to promote anotherdisplacive transformation to a rhombohedral structure preceding B19’ transformation, commonly known asR phase [1997Har]. Depending on the composition and thermal history of binary TiNi, and with the additionof Fe, the transformation temperatures B2 R B19’ may be well separated. The presence of the R-phaseis very useful for shape memory applications which rely on small thermal hysteresis.[1994Jia] reported the partitioning ratio of Fe, defined by xFe(TiNi3)/xFe(Ni), where xFe is the mole fractionof Fe, between (Ni) and TiNi3. The partitioning ratios at 1000, 1100 and 1200°C were 0.41, 0.44 and 0.61,respectively.There is no ternary phase in this system. The details of the crystal structures and lattice parameters of thesolid phases are listed in Table 1.


Even though there are no true pseudobinary sections, two sections have been reported as pseudobinary.[1938Vog] reported the TiFe2-TiNi3 section with the eutectic reaction L TiFe2+TiNi3 at 1320°C.[1967Dud] established the section TiFe-TiNi, which is shown in Fig. 1. The continuous solid solubilitybetween TiFe and TiNi was confirmed by lattice parameter and hardness measurements. [1967Dud]established only the solidus boundary. However, in the TiFe-end both liquid+TiFe2 andliquid+TiFe2+Ti(Fe,Ni) phase regions should appear due to formation of TiFe. The liquidus line expectedby [1967Dud] appears to be inconsistent with the investigated solidus. Also, in agreement with thecomments by [1991Gup] it is here reported with a minimum so that B2-Ti(Fe,Ni) melts congruently around1270°C.


Figure 2 shows the reaction scheme involving five invariant (U1, U2, E1, U3, U4) and a maximum (e1)reactions. Among these, e1, U1 and E1 were reported by [1938Vog] while U3 and U4 were reported by[1994Ali1]. The invariant reaction U2 has not been experimentally verified, but it was speculated by[1991Gup]. The composition of the phases participating in e1, U1, E1 and U3 are listed in Table 2. Thesecompositions were read from the superimposed liquidus and projection diagrams provided by [1938Vog],[1994Ali1] and [2001Gup].

Liquidus Surface

The liquidus surface and projection diagram for the composition range Fe-TiFe2-TiNi3-Ni was presentedby [1938Vog]. The superimposed liquidus surface and projection diagram for the Ti corner was reported by

301


MSIT®

Fe–Ni–Ti

[1994Ali1], and slightly modified by [2001Gup]. Using these results and the accepted binary phasediagrams, the liquidus surface shown in Fig. 3 was constructed.

Isothermal Sections

[1963Spe] reported partial isothermal sections of the Fe corner at 700 and 1100°C; however, did not providedetails of the experimental techniques and procedures. Figure 4 shows a partial isothermal section of theFe-rich corner at 1100°C [1963Spe]. Figure 5 and 6 show isothermal sections at 1000 [1999Abr] and 900°C[1981Loo], respectively. The results of [1981Loo] and [1999Abr] agree very well; however, a majordiscrepancy is that [1999Abr] observed very little solubility of Fe in Ti2Ni while [1981Loo] reported thatabout 78 % Ni sites in Ti2Ni can be substituted by Fe i.e, the solid solubility is about 26 at.% Fe. In anotherinvestigation, [1994Ali2] reported that only about 1.5 at.% Fe dissolves in Ti2Ni where rapidly solidifiedFe-Ni-Ti alloys were annealed at 900°C for only 25 h compared to up to 900 h by [1981Loo]. It is not clearif a short annealing treatment used by [1994Ali2] is responsible for the much lower solubility of Fe in Ti2Nicompared to [1981Loo]. The continuous solubility of TiFe and TiNi observed by [1981Loo] and [1999Abr]confirms the earlier results of [1967Dud]. In Figs. 5 and 6, the widths of TiNi and TiNi3 fields have beenadjusted to make them consistent with the accepted Ni-Ti phase diagram. The solubility of Fe in TiNi3 at900°C is about 15 at.% [1981Loo] which is much higher than reported by [1938Vog]. Figure 7 shows apartial isothermal section of the Fe corner at 700°C [1963Spe].


Several temperature-composition sections have been reported. Figures 8, 9 and 10 show polythermalsections at 8, 12 and 14.4 mass% Ti [1938Vog], respectively. Figures 11 and 12 showtemperature-composition sections at constant Fe:Ni mass ratios of 90:10 and 40:60, respectively[1938Vog]. The existence of two invariant reactions at 1200°C (U1) and at 1120°C (E1) is reflected in Figs.8 to 10. The partial isopleths of the Ti corner and at constant Fe:Ni atomic ratios of 1:3, 1:1 and 3:1 areshown in Figs. 13, 14 and 15, respectively [1994Ali1]. On the basis of these three isopleths, [1994Ali1] gavea superimposed partial liquidus surface and projection diagram for the Ti corner. Due to the existence of( Ti)+Ti(Fe,Ni) phase field in Fig. 13, the original projection diagram was slightly modified by [2001Gup].

Thermodynamics

[1975Ost] reported the enthalpy of dissolution of Ti in Fe-Ni melts. [1991Lue] measured the enthalpy ofmixing of liquid Tix(Fe0.89Ni0.11)1-x, 0.295 x 0.041, at 1600°C using a high-vacuum high-temperaturecalorimeter. They also modeled the molar heat of mixing of the entire Fe-Ni-Ti system using variousextrapolation methods. Subsequently, these experimental data were used to validate “thermodynamicadapted” power series for the extrapolation of thermodynamic quantities [1995Tom]. [1999Thi] also carriedout calorimetric measurement of heat of mixing of liquid alloys in three composition ranges: (i)(Fe84Ti16)1-xNix, 0.02 < x < 0.35 at 1621°C, (ii) (TiFe2)1-xNix, 0.02 < x < 0.4 at 1643°C, and (iii)(TiNi)1-xFex, 0.02 < x < 0.8 at 1645°C. [1999Thi] used an associate model to describe the heat of mixing ofternary alloys. Two associates, TiNi3 and TiFe, were assumed to be present in the liquid. In the entirecomposition range, the model calculations agreed very well with the experimental data implying that binaryinteractions were sufficient to describe the excess heat of mixing. Their model calculations also agreed verywell with the experimental data of [1991Lue]. The functional representation of experimental heat of mixingalong four composition sections is provided in Table 3. [1998Mie] performed CALPHAD modeling to calculate phase equilibria involving liquid, ( Fe) and ( Fe)phases. He introduced asymmetric ternary interaction parameters for the liquid phase and a symmetricternary interaction parameter for the bcc phase, but no ternary interaction parameter for the fcc phase. Thecalculated isothermal sections of Fe corner at 1100 and 1200°C were in good agreement with theexperimental data.

302


MSIT®

Fe–Ni–Ti


[1999Lee] studied hydriding behavior of TiFe1-xNix for x=0.1, 0.15 and 0.2. They found that partialsubstitution of Fe by Ni in TiFe does not change the hydriding behavior, except that these alloys can behydrided without the activation treatment.Ti0.5(Fe1-xNix)0.5 alloys are known to exhibit shape memory effect. [2000Xu] obtained a maximum shaperecovery strain of 5.6% after deforming Ti50Fe2Ni48 alloy at -70°C.

Miscellaneous

Discontinuous precipitation in ternary metastable alloys has been studied several times [1963Spe,1973Zem, 1979Fou, 1979Zem]. [1963Spe] investigated an 5.8Ti-Fe-29.7Ni (mass%) alloy in thetemperature range of 400 to 975°C, and [1979Fou] investigated an 6.5Ti-Fe-28.5Ni (mass%) alloy in thetemperature range of 400 to 900°C. [1979Zem] used (2 to 5)Ti-Fe-(25 to 26)Ni (mass%) alloys heat treatedat 790°C. According to [1963Spe], ternary alloys may exhibit two types discontinuous precipitations

1 + TiNi31 2 + (Fe,Ni)2Ti,

where 1 and 2 are austenitic solid solutions with different solute contents. Both [1979Fou] and[1979Zem] observed only the first discontinuous reaction where TiNi3 is a metastable phase because thealloys lie in the +TiFe2 phase field. The second discontinuous reaction is very sluggish [1963Spe].[1988Sag, 1992Sag] studied the effect of severe plastic deformation of austenitic 2.6Ti-Fe-36Ni alloycontaining equilibrium TiNi3 ( , DO24) and metastable TiNi3 ( ’, L12) precipitates which were obtained byadjusting the aging treatments. Upon severe plastic deformation both types of precipitates dissolve in thematrix due to strong interaction with dislocations. The effect of Fe on the martensitic transformations (B2 R B19’) in TiNi has been studied extensively[1985Goo, 1986Edm, 1990Rao, 1992Shi, 1993Mat, 1995Gue, 1997Har, 2000Chu, 2000Vor, 2000Xu].Addition of Fe decreases both martensitic and pre-martensitic transformation temperatures, and stabilizesthe R phase (rhombohedral). The martensitic transformation ( ) start temperature (Ms) ofFe-22.5Ni (mass%) [1963Yeo], Fe-27Ni (mass%) [1969Abr] and Fe-29.5Ni (mass%) [1969Abr] alloysdecreases with the addition of Ti.[1999Efi] reported the interdiffusion coefficient of ternary solid solutions ( ) at 1200°C with nickelcontents of 87 to 99 at.%. [1985Val] found that the average magnetic moment of ’-FeNi3 decreases when Fe is replaced by Ti.

References

[1938Vog] Vogel, R., Wallbaum, H.J., “The System Fe-Ni3Ti-Fe2Ti” (in German), Arch.

Eisenhuettenwes., 12, 299-304 (1938) (Equi. Diagram, Experimental, #, *, 16)[1941Vog] Vogel, R., “Über eine Beobachtung von Erzwunger Ausscheidungsrichtung in

Mischkristallen” (in German), Z. Metallkd., 33, 376-377 (1941) (Experimental, Equi.Diagram, 1)

[1963Spe] Speich, G. R., “Cellular Precipitation in an Austenitic Fe-30Ni-6Ti Alloy”, Trans. Met. Soc.

AIME, 227, 754-762 (1963) (Crys. Structure, Experimental, Equi. Diagram, #, *, 31)[1963Yeo] Yeo, R.G.B., “The Effects of Some Alloying Elements on the Transformation of

Fe-22.5%Ni Alloys”, Trans. Met. Soc. AIME, 227, 884-890 (1963) (Experimental, 13)[1967Dud] Dudkina, L.P., Kornilov, I.I., “An Investigation of the Equilibrium Diagram of the

TiNi-TiFe System”, Russ. Metall., (4), 98-101 (1967) (Experimental, Equi. Diagram,#, *, 13)

[1968Abr] Abraham, J.K., “Lattice Parameters as a Function of Ti in Austenitic Fe-Ni-Ti Alloys”,Trans. Met. Soc. AIME, 242, 2365-2369 (1968) (Crys. Strucutre, Experimental, 10)

[1969Abr] Abraham, J.K., Pascover, J.S., “The Transformation and Structure of Fe-Ni-Ti Alloys”,Trans. Met. Soc. AIME, 245, 759-768 (1969) (Crys. Structure, Experiment, 13)

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Fe–Ni–Ti

[1973Zem] Zemtsova, N.D., Malyshev, K.A., “Continuous Decomposition of a Gamma-Solid Solutionin Iron-Nickel-Titanium Alloys”, Phys. Met. Metallogr. (Engl. Transl.), 35(5), 108-114(1973) (Crys. Structure, Experimental, Equi. Diagram, 13)

[1975Ost] Ostrovskii, O.I., Dyubanov, V.G., Stomakhin, A.Ya., Grigoryan, V.A., “Enthalpy ofDissolution of Titanium in Iron-Nickel Melts”, Izvest. Vyssh. Ucheb. Zaved., Chern.

Metall., (7), 67-70 (1975) (Experimental, Thermodyn., 7)[1979Fou] Fournelle, R.A., “Discontinuous Coarsening of Lamellar Cellular Precipitates in an

Austenitic Fe-30Ni-6Ti Alloy”, Acta Metall., 27, 1135-1145 (1979) (Experimental, 18)[1979Zem] Zemtsova, N.D., Mabyshev, K.A., Starchenko, Ye. I., “Cellular Decomposition in

Undeformed Metastable Fe-Ni-Ti Alloys”, Phys. Met. Metallogr., 48(2), 128-135 (1979)(Experimental, Thermodyn., 11)

[1981Loo] van Loo, F.J.J., Vroljik, J.W.G.A., Bastin, G.F., “Phase Relations and Diffusion Paths in theTi-Ni-Fe System at 900°C”, J. Less-Common Met., 77, 121-130 (1981) (Experimental,Equi. Diagram, #, *, 11)

[1982Hwa1] Hwang, C.M., Meichle, M., Salmon, M.B., Wayman, C.M., “Transformation Behavior ofTi50Ni47Fe3 Alloy: I. Incommensurate and Commensurate Phases”, J. Phys., 43(12),C4-231-C4-236 (1982) (Crys. Structure, Experimental, 22)

[1982Hwa2] Hwang, C.M., Salmon, M.B., Wayman, C.M., “Transformation Behavior of Ti50Ni47Fe3Alloy: I. Martensitic Transformation”, J. Phys., 43(12), C4-237-C4-242 (1982) (Crys.Structure, Experimental, 5)

[1982Nis] Nishida, M., Honma, T., “Phase Transformations in Ti50Ni50-xFex Alloys”, J. Phys., 43(12),C4-225-C4-230 (1982) (Crys. Structure, Experimental, 6)

[1985Goo] Goo, E., Sinclair, R., “The B2 to R Transformation in Ti50Ni47Fe3 and Ti49.5Ni50.5”, Acta

Met., 33(9), 1717-1723 (1985) (Crys. Structure, Experimental, 29)[1985Gup] Gupta, K.P., Rajendraprasad, S.B., Jena, A.K., Sharma, R.C., “The Iron-Nickel-Titanium

System”, J.Alloy Phase Diagrams, 1, 59-68 (1985) (Equi. Diagram, Review, #, *, 21)[1985Val] Valiyev, E.Z., Men’shikov, A.Z., Panakhov, T.M., “The Structural State of Alloys

Ni3(Fe1- Tix) and Ni3(Fe1-xNbx) During Atomic Ordering”, Phys. Met. Metallogr., 59(1),123-129 (1985) (Crys. Structure, Experimental, 14)

[1986Edm] Edmonds, K.R., Hwang, C.M., “Phase Transformations in Ternary TiNiX Alloys”, Scr.

Metall., 20(5), 733-737 (1986) (Experimental, 7) [1988Sag] Sagaradze, V.V., Makozov, S.V., “Dissolution of Spherical and Platelike Intermetallics in

Fe-Ni-Ti Austenitic Alloys During Cold Plastic Deformation”, Phys. Met. Metallogr.,66(2), 111-120 (1988) (Crys. Structure, Experimental, 19)

[1990Rao] Rao, J., He, Y., Ma, R., “Study of Phase Transition in Ti50Ni47.5Fe2.5 Alloy”, Metall.Trans.,21A, 1322-1324 (1990) (Crys. Structure, 6)

[1991Gup] Gupta, K.P., “The Fe-Ni-Ti (Iron-Nickel-Titanium) System”, in “Phase Diagrams of

Ternary Nickel Alloys”, Indian Institute of Metals, Calcutta, India, 321-343 (1991) (Equi.Diagram, Review, #, *, 12)

[1991Lue] Lueck, R., Wang, H., Predel, B., “Thermodynamic Investigation of LiquidIron-Nickel-Titanium Alloys” (in German), Z. Metallkd., 82, 805-809 (1991)(Experimental, Thermodyn., 17)

[1991Mur] Murray, J.L., “Fe-Ti (Iron-Titanium)”, in “Phase Diagrams of Binary Iron Alloys”, ASMInternational, Metals Park, Ohio, 414-425 (1991) (Equi. Diagram, Review, #, *, 112)

[1992Sag] Sagaradze, V.V., Morozov, S.V., “Mechanical Alloying of Fe-Ni-Ti Austenite DuringLow-Temperature Deformation”, Mater. Sci. Forum, 88-90, 147-154 (1992) (Crys.Structure, Experimental, Equi. Diagram, 3)

[1992Shi] Shimizu, K., Tadaki, T., “Recent Studies on the Precise Crystal-Structural Analyses ofMartensitic Transformation”, Mater. Trans., JIM, 33(3), 165-177 (1992) (Crys.Structure, 101)

[1993Mat] Matveeva, N.M., Klopotov, A.A., Kormin, N.M., Sazanov, Yu.A., “Lattice Parameters andthe Sequence of Transformations in Ternary TiNi-TiMe Alloys”, Russ. Metall. (Engl.

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Fe–Ni–Ti

Transl.), (3), 216-220 (1993), translated from Izv. RAN, (3), 233-237 (1993)(Experimental, 10)

[1994Ali1] Alisova, S.P., Budberg, P.B., Barmina, T.I., Lutskaya, N.V., “Ti-TiFe-TiNi System”, Russ.

Metall. (Engl. Transl.), (1), 117-121 (1994), translated from Izv. RAN, (1), 158-163 (1994)(Experimental, Equi. Diagram, 9)

[1994Ali2] Alisova, S.P., Kovneristyi, Yu.K., Lutskaya, N.V., Budberg, P.B., “Structure of the RapidlySolidified Ti-TiFe-TiNi Alloys”, Russ. Metall. (Engl. Transl.), 146-149 (1994), translatedfrom Izv. RAN Met., (1), 158-161, 1994 (Experimental, Equi. Diagram, 5)

[1994Jia] Jia, C.C., Ishida, K., Nishizawa, T., “Partitioning of Alloying Elements Between (A1) and(DO24) Phases in The Ni-Ti Base Systems”, in “Experimental Methods Phase Diagram

Determination”, Morral, J.E., Schiffman, R.S., Merchant, S.M., (Eds.), The Minerals,Metals and Materials Society, Warrendale, PA, 31-38 (1994) (Experimental, Equi.Diagram, 8)

[1995Gue] Guerin, G., “Martensitic Transformation and Thermomechanical Properties”, Key Eng.

Mater., 101-102, 339-392 (1995) (Crys. Structure, Phys. Prop., Review, 73) [1995Tom] Tomiska, J., Wang, H., “On the Algebraic Evaluation of the Ternary Molar Heat of Mixing

HE from Calorimetric Investigations”, Ber. Bunsen-Ges. Phys. Chem., 99(4), 633-640(1995) (Thermodyn., 27)

[1997Har] Hara, T., Ohba, T., Okunishi, E., Otsuka K., “Structural Study of R-Phase in Ti-50.23Ni(at.%) and Ti-47.75Ni-1.50Fe (at.%) Alloys”, Mater. Trans., JIM, 38(1), 11-17 (1997)(Crys. Structure, Experimental, 18)

[1998Mie] Miettinen, J., “Approximate Thermodynamic Solution Phase Data for Steels”, Calphad,22(2), 275-300 (1998) (Equi. Diagram, Thermodyn., 83)

[1999Abr] Abramycheva, N.L., V'yunitskii, I.V, Kalmykov, K.B., Dunaev, S.F., “Isothermal CrossSection of the Phase Diagram of the Fe-Ni-Ti System.AT 1273 K. - I.”, Vestn. Mosk. Univ.,

Ser.2: Khim, 40(2), 139-143 (1999). (Experimental, Equi. Diagram, #, *, 4)[1999Efi] Efimenko, L.P., Petrova, L.P., Sviridov, S.I., “Interactions in TiFe-Ni System at 1200°C”,

Russ. Metall. (Engl. Transl.), (4), 160-165 (1999) (Experimental, 15) [1999Lee] Lee, S.M., Perng, T.P., “Correlation of Substitutional Solid Solution with Hydrogenation

Properties of TiFe1-xMx (M = Ni, Co, Al) Alloys”, J. Alloys Compd., 291, 254-261 (1999)(Crys. Structure, Experimental, 18)

[1999Thi] Thiedemann, U., Rösner-Kuhn, M., Drewes, K., Kuppermann, G., Frohberg, M.G.,“Temperature Dependence of the Mixing Enthalpy of Liquid Ti-Ni and Fe-Ti-Ni Alloys”,J. Non-Cryst. Solids, 250-252, 329-335 (1999) (Experimental, Thermodyn., 17)

[2000Chu] Chu, J.P., Lai, Y.W., Lin, T.N., Wang, S.F., “Deposition and Characterization of TiNi-BaseThin Films by Sputtering”, Mater. Sci. Eng. A, A277, 11-17 (2000) (Crys. Structure,Experimental, Phys. Prop., 20)

[2000Vor] Voronin, V.I., Naish, V.E., Novoselova, T.V., Sagaradze, I.V., “Structures of MonoclinicPhases in Titanium Nickelide: II. Transformation Cascade B2-R-T”, Phys. Met. Metallogr.,89(1), 19-26 (2000) (Crys. Structure, Experimental, 15)

[2000Xu] Xu, H., Jiang, C., Gong, S., Feng, G., “Martensitic Transformation of the Ti50Ni48Fe2 AlloyDeformed at Different Temperatures”, Mater. Sci. Eng. A, A281, 234-238 (2000)(Experimental, 11)

[2001Gup] Gupta, K.P., “The Fe-Ni-Ti System Update (Iron - Nickel - Titanium)”, J. Phase Equilib.,22(2), 171-175 (2001) (Equi. Diagram, Review, #, *, 5)

[2004Kuz] Kuznetsov, V, “Fe-Ni (Iron-Nickel)”, MSIT Binary Evaluation Program, in MSIT

Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH,Stuttgart; to be published, (2004) (Crys. Structure, Equi. Diagram, Assessment, 41)

[2004Ted] Tedenac, J.C., Velikanova, T., Turchanin, M., “Ni-Ti (Nickel-Titanium)”, MSIT BinaryEvaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), Materials ScienceInternational Services, GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure,Equi. Diagram, Assessment, 37)

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Phase/ Temperature Range [°C]

Pearson Symbol/ Group Space/Prototype

Lattice Parameters [pm]

Comments/References

( Fe)(h2) 1538 - 1394

cI2Im3m

W

a = 293.15 pure Fe at 1390°C [Mas2]

, ( Fe,Ni)

( Fe)(h1) 1394 - 912 (Ni) < 1455

cF4Fm3m

Cu

a = 358.37a = 359.13a = 358.90a = 359.39a = 359.0a = 365.2a = 364.67

a = 352.32

Fe71.74Ni27.07Ti1.19, at 20°C [1968Abr]Fe68.27Ni26.96Ti4.77, at 20°C [1968Abr]Fe68.32Ni30.38Ti1.3, at 20°C [1968Abr]Fe64.39Ni29.65Ti5.96, at 20°C [1968Abr]Fe31Ni63Ti6, at 20°C [1981Loo] Fe31Ni63Ti6, at 900°C [1981Loo]pure Fe [Mas2]

pure Ni at 20°C [V-C2]

( Fe)(r) < 912

cI2Im3m

W

a = 286.65 pure Fe at 20°C [V-C2]

( Fe) hP2P63/mmc

Mg

a = 246.8c = 396.0

at 25°C, 13 GPa [V-C2]

( Ti)(h) 1670 - 882

cI2Im3m

W

a = 330.65 [Mas2]

( Ti)(r) 882

hP2P63/mmc

Mg

a = 295.06 c = 468.25


'-FeNi3 517

cP4Pm3m

AuCu3

a = 355.23 63 to 85 at.% Ni [2004Kuz]

TiFe2 1427

hP12P63/mmc

MgZn2

a = 478.7 c = 781.5

24.0 to 36.0 at.% Ti [V-C2]

dissolves up to 28 at.% Ni [1981Loo]

Ti(Fe,Ni)

TiFe 1317

TiNi 1311

cP2Pm3m

CsCl

a = 300.0a = 298.91a = 298.18a = 297.6

a = 299.8 to 301.0

Fe10.2Ni39.8Ti50, at 20°C [1967Dud]Fe25.3Ni24.7Ti50, at 20°C [1967Dud]Fe35.3Ni14.7Ti50, at 20°C [1967Dud]49.8 to 51.8 at.% Ti [V-C2]

49.5 to 57 at.% Ni [2004Ted]

Ti2Ni 984

cF96Fd3m

Ti2Ni

a = 1127.8 to 1132.4

33 to 34 at.% Ni [2004Ted]

dissolves up to 26 at.% Fe [1981Loo]

306


MSIT®

Fe–Ni–Ti

Table 2: Composition of the Phases Participating in the Reaction Scheme of Fe-Ni-Ti System

TiNi (martensite) hP18P3

-

a = 735.8 c = 528.55

Ni50.23Ti49.77, at 20°C [1997Har].X-ray diffraction. Known as R phase.

TiNi (martensite) mP4P21/m

TiNi

a = 289.8 b = 410.8c = 464.6

= 97.78°

Ni49.2Ti50.8, at 20°C [1992Shi].Single crystal X-ray diffraction. Known as B19’ martensite.

TiNi3 1380

hP16P63/mmc

TiNi3

a = 510.28c = 827.19

a = 510.3 0.5c = 832.0 0.8a = 517.0c = 846.8

75 to 80.1 at.% Ni at 1300°C [2004Ted]

dissolves up to 14 at.% Fe [1981Loo]Fe4Ni72Ti24, at 20°C [1981Loo]Fe4Ni72Ti24, at 900°C [1981Loo]

Reactions T [°C] Type Phase Composition (at.%)

Fe Ni Ti

L TiFe2 + TiNi3 1320 e1 LTiFe2TiNi3

31.80 52.46 0.50

39.10 16.61 73.26

29.10 30.93 26.24

L + ( Fe) ( Fe) + TiFe2 1200 U1 L( Fe)( Fe)TiFe2

66.46 79.40 76.94 59.71

17.55 11.33 13.90 8.21

15.99 9.269.1632.01

L ( Fe) + TiFe2 + TiNi3 1113 E1 L( Fe)TiFe2TiNi3

36.35 55.29 52.46 0.50

40.61 35.68 16.61 73.26

23.04 9.0330.9326.24

L + Ti(Fe,Ni) ( Ti) + Ti2Ni 960 U3 LTi(Fe,Ni)( Ti) Ti2Ni

0.91 18.02 0.95 1.11

24.03 28.57 4.04 32.64

75.0653.4195.0166.25


Pearson Symbol/ Group Space/Prototype

Lattice Parameters [pm]

Comments/References

307


MSIT®

Fe–Ni–Ti

Table 3: Heat of Mixing ( Hmix) of Fe-Ni-Ti Liquid Alloys. The Data Below are Due to Functional Representation of Experimental Data [1999Thi].

Composition Section Temperature [°C] Hmix, [kJ (g-at.)-1]

(Fe84Ti16)1-yNiy 1621

y = 0.0y = 0.05y = 0.10y = 0.15y = 0.20y = 0.25y = 0.30y = 0.35

-9.291-9.843-10.433-10.787-11.024-11.083-11.024-10.945

(NiTi)1-yFey 1645

y = 0.0y = 0.05y = 0.10y = 0.15y = 0.20y = 0.25y = 0.30y = 0.35

-36.972-34.507-32.141-30.169-28.197-26.127-24.648-23.070

(Fe2Ti)1-yNiy 1643

y = 0.0y = 0.05y = 0.10y = 0.15y = 0.20y = 0.25y = 0.30y = 0.35

-17.362-17.872-18.191-18.511-18.511-18.415-18.670-17.681

(Fe89Ni11)1-yTiy 1600

y = 0.0y = 0.05y = 0.10y = 0.15y = 0.20y = 0.25y = 0.30y = 0.35

-1.892-5.081-8.108-10.811-13.514-15.892-17.946-20.004

308


MSIT®

Fe–Ni–Ti

10 20 30 40

1100

1200

1300

1400

1500

Ti 50.00

Ni 50.00

Fe 0.00

Ti 50.00

Ni 0.00

Fe 50.00Fe, at.%

Te

mp

era

ture

, °C

L+TiFe2+Ti(Fe,Ni)

L+TiFe2

L

L+Ti(Fe,Ni)

Ti(Fe,Ni)

Fig. 1: Fe-Ni-Ti.

The TiNi-FeTi pseudobinary section

309


MSIT®

Fe–Ni–Ti

Fig

. 2:

F

e-N

i-T

i. R

eact

ion s

chem

e

L +

(δF

e)γ

15

17

p1

Lγ

+ T

iNi 3

13

04

e 2

L T

iFe 2

+ T

iNi 3

13

20

e 1

γ 1(α

Fe)

+γ 2

41

2e 9

γ 2(α

Fe)

+ F

eNi 3

34

5e 1

0

L T

i(F

e,N

i)+

TiN

i 3

11

20

e 4

L+

Ti(

Fe,

Ni)

TiN

i

98

5p

3

L (

βTi)

+ T

i 2N

i

94

2e 6

(βT

i) (

αTi)

+T

i 2N

i

76

7e 7

L+

TiF

e 2T

i(F

e,N

i)

13

17

p2

L(α

Fe)

+T

iFe 2

12

90

e 3

L(β

Ti)

+T

i(F

e,N

i)

10

85

e 5

(βT

i)(α

Ti)

+T

i(F

e,N

i)

59

0e 8

Ni-

Ti

Fe-

Ni

Fe-

Ti

Fe-

Ni-

Ti

L+

γ+T

iFe 2

(δF

e)+

γ+T

iFe 2

TiF

e 2+

Ti(F

e,N

i)+

TiN

i 3L

+T

i(Fe,

Ni)

+T

iNi 3

γ+T

iFe 2

+T

iNi 3

L+

(βT

i)+

Ti 2

Ni

Ti(

Fe,

Ni)

+(β

Ti)

+T

i 2N

i

(αT

i)+

(βT

i)+

Ti(

Fe,

Ni)

(αT

i)+

Ti(

Fe,

Ni)

+T

i 2N

i

L +

(δF

e)γ

+ T

iFe 2

12

00

U1

L+

TiF

e 2T

i(F

e,N

i)+

TiN

i 3?

U2

Lγ

+ T

iFe 2

+ T

iNi 3

11

13

E1

L+

Ti(

Fe,

Ni)

(βT

i)+

Ti 2

Ni

96

0U

3

(βT

i)+

Ti 2

Ni

(αT

i)+

Ti(

Fe,

Ni)

65

0U

4

310


MSIT®

Fe–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Fe


Axes: at.%

(βTi)

Ti(Fe,Ni)TiFe

2

TiNi3

(δFe)

γ

U3

p3

e5

e4

e2

e6

p2

e3

p1

E1

U2

e1

U1

NiTi2

10

20

30

40

60 70 80 90

10

20

30

40

Ti 50.00

Fe 50.00

Ni 0.00

Fe

Ti 0.00

Fe 50.00

Ni 50.00Data / Grid: at.%

Axes: at.%

γ+TiN

i 3+T

iFe 2

γ+TiNi3

γ

γ+TiFe2

(αFe)

(αFe)+γ(αFe)+γ+TiFe2

(αFe)+TiFe2

Fig. 3: Fe-Ni-Ti.

Liquidus surface

Fig. 4: Fe-Ni-Ti.

Partial isothermal section at 1100°C

311


MSIT®

Fe–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Fe


Axes: at.%

Ti2Ni+(βTi)

+TiNi3

γ+TiFe2

γ

(βTi)

+TiFe2

(αFe)

(Ni)+TiNi3

TiNi3+Ti(Fe,Ni)

Ti(Fe,Ni)+

+Ti2Ni

Ti(Fe,Ni)+ γ+TiFe2+

TiNi3+

+Ti(Fe,Ni)+Ti2Ni

+TiFe2

Ti(Fe,Ni)

TiNi3

TiFe2

20

40

60

80

20 40 60 80

20

40

60

80

Ti Fe


Axes: at.%

γ

Ti(Fe,Ni)

TiFe2

Ti2Ni

L

(βTi)

(αFe)

TiNi3

Fig. 6: Fe-Ni-Ti.

Isothermal section at 900°C

Fig. 5: Fe-Ni-Ti.


312


MSIT®

Fe–Ni–Ti

10

20

30

40

60 70 80 90

10

20

30

40

Ti 50.00

Fe 50.00

Ni 0.00

Fe

Ti 0.00

Fe 50.00

Ni 50.00Data / Grid: mass%

Axes: mass%

(αFe)+γ+TiFe2

(αFe)

γ

(αFe)+γ

γ+TiNi3

γ+T

iNi 3

+T

iFe 2

(αFe)+TiFe2

γ+TiFe2

20 40 60 80

500

750

1000

1250

1500

Ti 9.21

Ni 0.00

Fe 90.79

Ti 9.63

Ni 90.37

Fe 0.00Ni, at.%

Te

mp

era

ture

, °C

γ

γ+TiNi3

L+γ

γ+TiFe2+TiNi3

L+γ+TiNi3L+γ+TiFe2

(αFe)

L+(αFe)+γ

L+(αFe)

(αFe)+TiFe2

γ+TiFe2

(αFe)+γ

(αFe)+γ+TiFe2

Fig. 7: Fe-Ni-Ti.


Fig. 8: Fe-Ni-Ti.

A polythermal section at a constant Ti content of 8 mass%

313


MSIT®

Fe–Ni–Ti

20 40 60 80

500

750

1000

1250

1500

Ti 13.72

Ni 0.00

Fe 86.28

Ti 14.32

Ni 85.68

Fe 0.00Ni, at.%

Te

mp

era

ture

, °C

γ+TiNi3

L+γ+TiNi3

(αFe)+γ+TiFe2

γ+TiFe2

L+(αFe)+TiFe2

L+γ

γ+TiFe2+TiNi3

L+γ+TiFe2

L+(αFe)

L+(αFe)+γ

(αFe)+

TiFe2

L

γ

20 40 60 80

500

750

1000

1250

1500

Ti 16.40

Ni 0.00

Fe 83.60

Ti 17.10

Ni 82.90

Fe 0.00Ni, at.%

Te

mp

era

ture

, °C

γ+TiNi3

L+γ+TiFe2

L+γ+TiNi3

γ+TiFe2

L+γ

γ+TiFe2+TiNi3

(αFe)+γ+TiFe2

L+(αFe)

L+γ+TiFe2

(αFe)+

TiFe2

L+(αFe)+TiFe2

Fig. 9: Fe-Ni-Ti.

A polythermal section at a constant Ti content of 12 mass%

Fig. 10: Fe-Ni-Ti.

A polythermal section at a constant Ti content of 14.4 mass%

314


MSIT®

Fe–Ni–Ti

10 20 30

500

750

1000

1250

1500

Ti 0.00

Ni 9.56

Fe 90.44

Ti 33.37

Ni 2.27

Fe 64.36Ti, at.%

Te

mp

era

ture

, °C

L+TiFe2

(αFe)+γ

L+(αFe)(αFe)

(αFe)+γ+TiFe2

L+(αFe)+γ

γ

(αFe)+TiFe2

L+(αFe)+TiFe2

10 20

500

750

1000

1250

1500

Ti 0.00

Ni 38.82

Fe 61.18

Ti 28.59

Ni 41.99

Fe 29.42Ti, at.%

Te

mp

era

ture

, °C

γ

γ+TiNi3

L+γ

γ+TiNi3+TiFe2

L+TiNi3

L+γ+TiNi3L+TiNi3+TiFe2

L

Fig. 11: Fe-Ni-Ti.

A polythermal section at a constant mass ratio of Fe:Ni=90:10

Fig. 12: Fe-Ni-Ti.

A polythermal section at a constant mass ratio of Fe:Ni=60:40

315


MSIT®

Fe–Ni–Ti

10 20 30

500

600

700

800

900

1000

1100

1200

Ti Ti 50.00

Ni 37.50

Fe 12.50Ni, at.%

Te

mp

era

ture

, °C

L

Ti(Fe,Ni)

L+(βTi)+Ti(Fe,Ni)

(βTi)+Ti(Fe,Ni)

L+(βTi)

(βTi)

L+Ti(Fe,Ni)

L+Ti(Fe,Ni)+Ti2Ni

Ti(Fe,Ni)+Ti2Ni

(αTi)

(βTi)+Ti2Ni

(βTi)+Ti2Ni+Ti(Fe,Ni)(αTi)+(βTi)

(αTi)+Ti(Fe,Ni)+Ti2Ni

(αTi)+(βTi)+Ti2Ni

(αTi)+Ti2Ni

10 20

500

600

700

800

900

1000

1100

1200

Ti Ti 50.00

Ni 25.00

Fe 25.00Ni, at.%

Te

mp

era

ture

, °C

L

(βTi)+Ti(Fe,Ni)

Ti(Fe,Ni)

(βTi)

L+Ti(Fe,Ni)

L+(βTi)+Ti(Fe,Ni)

L+(βTi)

(αTi) (αTi)+(βTi)+Ti2Ni

(αTi)+(βTi)+Ti(Fe,Ni)

(αTi)+(βTi)

Fig. 13: Fe-Ni-Ti.

Partial polythermal section along the line of constant Fe:Ni atomic ratio of 1:3

Fig. 14: Fe-Ni-Ti.

Partial polythermal section along the line of constant Fe:Ni atomic ratio 1:1

316


MSIT®

Fe–Ni–Ti

90 80 70 60

500

600

700

800

900

1000

1100

1200

Ti Ti 50.00

Ni 12.50

Fe 37.50Ti, at.%

Te

mp

era

ture

, °C

L

Ti(Fe,Ni)

(βTi)+Ti(Fe,Ni)

L+(βTi) L+Ti(Fe,Ni)

L+(βTi)+Ti(Fe,Ni)

(βTi)

(αTi)

(αTi)+Ti(Fe,Ni)

(αTi)+(βTi)



Fig. 15: Fe-Ni-Ti.

Partial polythermal section along the line of constant Fe:Ni atomic ratio of 3:1

317


MSIT®

H–Mg–Ni

Hydrogen – Magnesium – Nickel

Pierre Perrot, Sigrid Wagner, Elena Semenova

Literature Data

The intermetallic compound Mg2Ni reacts readily with H2 at about 300°C under a moderate hydrogen

pressure of 3.28 bar. The product of the reaction is Mg2NiH4 [1968Rei, 1980Min]. If Mg is present in the

alloy, the pressure-composition isotherm exhibits two plateaus. The lower plateau (1.5 bar at 300°C) is due

to the known reaction (1):

Mg + H2 MgH2 (1)

When all the free Mg is exhausted, the second and higher plateau appears (3.28 bar at 300°C), which is due

to the reaction (2):

½Mg2Ni + H2 ½Mg2NiH4 (2)

Both reactions are reversible, the desorption peak occurring at a temperature lower than the absorption peak

owing to hysteresis. [1998Son] measured, between 275 and 320°C, and [2001Li], between 300 and 350°C,

the pressure for absorption of H by Mg2Ni and for desorption of H in Mg2NiH4. The best fit between both

measurements is given by:

RT ln (p/bar) = -57 860 + 112.1 T (absorption)

RT ln (p/bar) = -59 550 + 112.8 T (desorption)

With the deuteride, absorption and desorption isotherms are observed at slightly higher pressures

[1980Sch1].

The second intermetallic compound MgNi2 did not react at 300°C up to 2.7 MPa H2. However, a new

compound MgNi1.8H2 has been formed at 800°C under 5 MPa in an anvil type apparatus. It is thermally

stable up to 225°C under 0.5 MPa pressure of Ar [2000Kak].

A careful analysis of the thermogravimetric and differential thermal analysis curves [1988Sel] shows that

the decomposition of Mg2NiH4 follows a multi-step process, suggesting the existence of a new phase

Mg2NiH. However, [1998Kom] showed that the desorption peaks are strongly affected by air exposure after

hydrogenation. The shift of both peaks after air exposure varies between 50 and 100 K towards the higher

temperatures. [1999Zen] in a thermodynamic evaluation of the H-Mg-Ni ternary system does not take into

account the Mg2NiH phase.

The solubility of hydrogen in Mg and its alloys was measured by [1974Hua, 1976Wat] using a Sieverts

apparatus between 200 and 750°C. The hydrogen contents of the samples were determined by fusing the

alloys at about 1000°C and analyzing quantitatively the amount of hydrogen evolved. Nickel increases the

solubility of hydrogen in liquid magnesium. The intermetallic compound Mg2Ni can accommodate only 0.3

hydrogen atoms per formula unit. However, mechanical grinding of the intermetallic compound Mg2Ni

under a hydrogen atmosphere [1996Ori] leads to the composition Mg2NiH1.8 without changing the crystal

structure of the matrix Mg2Ni.

Binary Systems

The Mg-Ni system is accepted from [Mas2]. An evaluation by [1993Jac] included the magnetic behavior of

Ni and reproduced within 3°C the main features of the Mg-Ni phase diagram assessed by [1985Nay] and

accepted by [Mas2]. The H-Ni system has been assessed by [1989Way] who plotted the solubility of H in

solid Ni at 0.1 MPa pressure and the solubility of H in solid and liquid Ni between 1400 and 1460°C at 50

MPa pressure. [1987Sch] gives the solubility of H in liquid Ni at 0.1 MPa pressure between 1500 and

1700°C. A more recent evaluation [1999Zen] of the H-Ni system at 0.1 MPa bar pressure shows a retrograde

solubility of hydrogen in liquid nickel. The H-Mg system under a hydrogen pressure of 25 MPa is accepted

as given by [Mas2] from an assessment of [1987San]. The accepted H-Mg phase diagram at 0.1 MPa is

reported in [2001Per].

318


MSIT®

H–Mg–Ni

Solid Phases

The solid phases are presented in Table 1. In addition to Mg2NiH4 several compounds that look like ternary

hydrides have been reported in the literature: Mg2NiH0.3 [1982And, 2002Li], Mg2NiH [1988Sel] and

Mg2NiH1.8 [1996Ori]. Actually, these phases are solutions of hydrogen in Mg2Ni; the first one is stable; the

other two represent probably oversaturated, thus metastable, solutions of H into Mg2Ni. Mg2Ni3H3.4 is a

high pressure phase prepared from a mixture MgH2+Ni heated at 800°C under a pressure of 5 GPa obtained

with a cubic anvil type apparatus [2002Tak]. Mg2NiH4 undergoes several phase transformations under very

high pressures [2003Yam]. A first transformation monoclinic - cubic is induced at 1.3 - 2 GPa. A further

transformation to the Mg2NiH4(h2) cubic form takes place upon heating above 200°C at 5 GPa, followed,

above 300°C by the appearance of a new high pressure form, which may be the result of the

disproportionation reaction: Mg2NiH4 MgNiH2+MgH2. In presence of an internal hydrogen source

(LiAlH4 or NaBH4+Ca(OH)2), NiH forms as the result of the reaction: MgNiH2+H2 MgH2+2NiH

[2003Yam].

The only known stable ternary hydride is Mg2NiH4 which may be easily prepared by mechanical alloying

at room temperature under hydrogen atmosphere [2000Wan] or by hydriding combustion synthesis

[2000Li]. At temperatures above 506°C, the eutectic point of the Mg-Ni system, Mg2NiH0.3 forms as an

impurity which reduces the hydriding activity of Mg2Ni [2002Li]. The hydrogenation obtained by

combustion synthesis at lower temperatures produces pure Mg2NiH4 [2003Sai]. A mixture of Ni and Mg

powders needs only 1 MPa hydrogen pressure during the heating step from room temperature to 600°C,

which is much lower than 4 MPa needed when compact mixtures were used.

Three polymorphs were identified by X-ray and neutron diffraction [1984Dar, 1984Hay]. The room

temperature form, monoclinic with hydrogen distributed on a partially filled site [1984Nor], transforms

irreversibly into the Mg2NiH4(h2) cubic structure at 245°C under a pressure of 0.1 MPa and the

transformation gave a well defined peak in differential scanning calorimetry. In contrast, the orthorhombic

form Mg2NiH4(h1) transforms reversibly into the cubic Mg2NiH4(h2) structure at about 235°C and 0.1 MPa

[1984Dar] and can be prepared by cooling the h2 form to room temperature [1984Nor]. The orthorhombic

form Mg2NiH4(h1) is thus probably metastable. The three polymorphs can be described as a distortion of

the high temperature cubic cell Mg2NiH4(h2), [1984Hay, 1984Nor]. The parameters given by [1984Dar] for

the monoclinic form Mg2NiH4(r), obtained by neutron diffraction, are preferred to those given by

[1982Ono2] because they facilitate the comparison with the cubic form Mg2NiH4(h2).

[1984Nor] proposes two room temperature forms Mg2NiH4(r1) and Mg2NiH4(r2), the former being a

substructure of the latter, with a double a parameter, and showed that the diffraction pattern at room

temperature depends on the preparation of the sample. According to [1987Pos], calorimetric data

distinguish three low temperature and two high temperature modifications with an actual stoichiometry of

Mg2NiH3.91 0.03, independent of the temperature. Final observations of [1986Nor] and [1986Zol] showed

that only one room temperature modification exists. The low temperature modifications seem to be related

to a reorientational motion of the NiH4 clusters in the structure due to microtwins or stacking faults

introduced by grinding [1999Roe, 2002Blo]. The ordered monoclinic structure Mg2NiH4(r) is destroyed by

mechanical grinding [1999Roe] to form the disordered cubic structure Mg2NiH4(h2) (a = 649.23 nm) found

previously only above 245°C. The desorption of Mg2NiD4, occurs between 20 and 350°C without any phase

transformation [1981Sch]. Owing to the history of the sample different amounts of stacking faults can be

induced in the lattice.

The position of hydrogen atoms in the lattice had been investigated on the deuterated compounds

Mg2NiD4(h2) and Mg2NiD4(r) [1980Sch1, 1984Dar] by diffusion studies and nuclear magnetic resonance

[1984Sen, 1987Sen]. The thermal expansion of Mg2NiH4 between 0 and 400°C under a hydrogen pressure

of 20 bar has been measured by [1982Ono1].

The evaluation of relative stabilities of hypothetical compounds Mg2NiH ( = 2, 4, 6) agree with the

observation that Mg2NiH4 is the only stable compound [1999Gar]. The H distribution that minimizes total

energy corresponds to a tetrahedrally distorded coordination of Ni atoms with a bond length of 0.1548 nm

in good agreement with the neutron diffraction experimental data.

319


MSIT®

H–Mg–Ni


The pseudobinary section Mg2Ni-Mg2NiH4 presented in Fig. 1 is constructed from the isobaric curves of

[2000Kuj, 2001Kuj] with the experimental domains of [2002Kuj]. The hydrogen pressure at equilibrium

Mg2Ni-Mg2NiH4 is 0.33 MPa at 300°C, which agrees well with the hydrogen pressure of 0.328 MPa at

298.6°C calculated by [1999Zen]. The hydrogen solubility in solid Mg2Ni at 300°C corresponds to the

composition Mg2NiH0.3 (9.09 at.% H) given by [1968Rei, 1982And].

Isothermal Sections

Ni increases the solubility of H in (Mg). The solubility of hydrogen in liquid Mg-Ni alloys at 700°C, given

in Fig. 2 is calculated from the data of [1999Zen]. Solubilities calculated under 0.1 MPa are excellent

agreement with the experimental results of [1974Hua] in the same conditions (0.1 MPa, 700°C). The

solubility measurements of [1976Wat] have not been taken into account because they contradict with the

accepted binary Mg-Ni diagram: at constant temperature, in a two-phase region, the solubility of hydrogen

must be proportional to the atomic fraction of nickel in the alloy, which was not verified.

The isothermal section at 301°C under 3.47 bar is reported in Fig. 3 after the calculation of [1999Zen]. The

hydrogen solubility in the Mg2Ni phase is calculated to be 9.14 at.% whereas the experimental solubility

corresponding to the formula Mg2NiH0.3 is 9.09 at.%.

Thermodynamics

The enthalpy change associated to the transition Mg2NiH4(r) Mg2NiH4(h2) is evaluated at 6.7 kJ mol-1

[1980Min] and 7 kJ mol-1 [1985Nor]. From the slope of the plot of lnpeq vs 1/T, the enthalpy change of

reactions (1) and (2) were calculated by [1982Aki], respectively -74.5 and - 63.6 kJ mol-1 for the reaction

(2). [1983Sem] proposes respectively - 74.3 and - 64.4 kJ mol-1. A more recent determination of the

hydrogen pressure at equilibrium Mg2Ni-Mg2NiH4 leads [2000Kuj] to propose rH = - 62.6 kJ mol-1 for

the reaction (2) at high temperatures (T > 235°C). Unfortunately, the value proposed for the same reaction

at low temperatures ( rH = - 31.4 kJ mol-1) leads to an enthalpy change of 31.2 kJ mol–1 for the transition

Mg2NiH4(r) Mg2NiH4(h2), which is hardly credible. Even if this high value is also the value

(31.2 kJ mol-1) calculated ab initio by [2002Mye] from the total energy density functional theory, the most

probable enthalpy of transition, deduced from the isobaric temperature composition curves is trH = 7.4

0.9 kJ mol-1. The enthalpy of formation calculated from the extended Miedema’s model [2002Her] for

Mg2NiHx varies from - 55 kJ mol-1 for x = 3.73 to - 51 kJ mol-1 for x = 4. The most probable value is - 64

kJ mol-1 for Mg2NiH4(h2). The absorption isotherms were drawn by [2001Kuj] for different alloys Mg2-xNi

(x = 0, 0.25, 0.50). The transition temperature was found to decrease from 250 to 210°C when increasing

from x = 0 to x = 0.50 the nickel content of the alloy.

Thermodynamic quantities were investigated by emf studies between 142 and 170°C with an

electrochemical cell:

(-) Na/Na+ H- in NaAl(C2H5)4 // Alloy / Hydride (+).

The solid state phase diagram deduced is valid between these temperatures [1985Lue1, 1985Lue2,

1987Lue].

[1997Rud] proposes a model for the binary systems H-Mg, H-Ni and Mg-Ni, as well as for the liquid phase

of the H-Mg-Ni ternary system. [1998Rud] analyzed further the Mg-Ni system for its glass formation ability

and the influence of dissolved hydrogen upon the amorphization properties of the alloy.


Mg2Ni, as many other Laves phases, forms hydrides with large hydrogen capacity. However, many of these

hydrides decompose after few cycles of absorption/desorption, which shows that these hydrides are

metastable [1978Sha]. Mg2NiH4 had been proposed [1968Rei, 1976Sem, 1985Lue1, 1985Son, 1986Nor,

1988Sel, 1997Rud, 2000Aiz] for hydrogen storage because it releases hydrogen at a convenient temperature

following the reversible reaction: Mg2NiH4 Mg2Ni+2H2. Pure magnesium absorbs hydrogen gas to about

8.2% of its weight (3.7% for Mg2Ni), but the reaction rate between magnesium and hydrogen is too slow

320


MSIT®

H–Mg–Ni

[1982Aki, 2000Wan] and the stability of magnesium hydride is too low [1980Min], which precludes its use

for hydrogen storage purposes. Moreover, the volume content is higher as a result of an increase in the

density of the Mg2NiH4 phase as compared to the MgH2 phase [1976Sem]. Mg2Ni has been shown to

present a greater sorption/desorption rate and a weaker loss of hydrogen capacity than classical materials

such as Mg2Cu [1978Gui] or FeTi [1980Sch2]. The cycling stability of commercial crystalline Mg2Ni alloy

was tested over 2700 cycles of absorption/desorption at respective temperatures of 250 and 300°C by

[1999Deh]. During two cycles of heating then cooling in a temperature range from 25 to 550°C under a

hydrogen pressure of 1.0 MPa, [1998Li] put into evidence the following steps: hydriding, dehydriding, then

formation of Mg2Ni during the first heating period, formation of Mg2NiH4 during the first cooling period.

During the following cycles, Mg2NiH4 decomposes during heating, then forms during cooling. The

characteristics of the alloys, i.e. the thermodynamic properties, absorption/desorption rates, scanning

electron micrograph, specific surface area, heat capacity and crystal structure remains stable through 1500

cycles. A synergetic effect has been observed by [1999Zal] through ball-milling of the mixtures

Mg2NiH4-MgH2, allowing the mixture to operate at temperatures around 220-240°C, lower than

conventional MgH2, with excellent absorption/desorption kinetics and a hydrogen capacity exceeding 5%,

which proves that MgH2 participates in the reaction. This result is remarkable in that the dissociation of

MgH2 does not normally occur below 280°C [1999Zen].

Hydride films have attracted much attention owing to their potential uses in the domain of hydrogen storage.

Mg2Ni forms stable Mg2NiH4 films whereas MgNi2 does not react within 450°C and 4 MPa H2. Mg1.2Ni

films present reversible hydriding/dehydriding characteristics with moderate conditions (<205°C)

[2002Che].

Mg2NiH4 presents an irreversible conductor-insulator transition in the temperature interval -163 to +300°C

[2002Blo]. The disappearance of the electric conductivity is concomitant with the appearance of stacking

faults, or microtwinning in the structure. By compressing hydride samples, Mg2NiH4 regains its electric

conductivity as the observable amount of stacking fault is reduced.

Miscellaneous

Resistivity measurements showed that Mg2NiH4 is a semiconductor with an activation energy of 0.05 eV

(4.8 kJ mol-1), which reflects the covalent character of the hydrogen-nickel bonding in the NiH4 complex

[1985Nor]. Energy [1989Hua1] and entropy [1989Hua2] considerations suggest that, in cubic Mg2NiH4 at

a temperature range around 300-400°C, nearly two thirds of the H atoms are in a square planar configuration

around Ni atoms and one third in distorted tetrahedral configurations around Ni atoms. In the hydriding

process of Mg2Ni, the rate-controlling step is the dissociative chemisorption of hydrogen [1985Son].

Infrared, Raman and Inelastic neutron scattering spectra of Mg2NiH4 have been observed and assigned by

[2002Par].

The hydrogen storage characteristics of Mg-Ni alloys may be improved by the addition of a third element

such as Nd [2001Yin] or graphite [2004Bob]. For instance, Mg86Ni10Nd4 can absorb 5 mass% H at

excellent speed between 100 and 300°C under 3 MPa H2 and can desorb at moderate speed between 300

and 200°C. Graphite was shown to improve the hydriding characteristics of Mg2Ni at 300°C and below.

The composite 90 mass% Mg2Ni+10 mass% graphite has a two times quicker desorption times than the non

modified Mg2Ni.

References

[1968Rei] Reilly, J.J., Wiswall, R.H., “The Reaction of H with Alloys of Mg and Ni and the Formation

of Mg2NiH4”, Inorg. Chem., 7, 2254-2256 (1968) (Thermodyn., #, 4)

[1974Hua] Huang, Y.C., Watanabe, T., Komatsu, R., “Hydrogen in Magnesium and its Alloys”, Proc.

Internat. Conf. Vacuum Metallurgy, 4th, 1973, 176-179 (1974) (Experimental, 8)

[1976Sem] Semenenko, K.N., Burnashova, V.V, “Interaction of Intermetallic Compounds with

Hydrogen”, Dokl. Chem., 678-680, (1976), translated from Dokl. Akad. Nauk SSSR, 231,

356-358 (1976) (Experimental, 12)

321


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H–Mg–Ni

[1976Wat] Watanabe, T., Tachihara, Y., Huang, Y.C., Komatsu, R., “The Effect of Various Alloying

Elements on the Solubility of Hydrogen in Magnesium”(in Japanese), J. Jpn. Inst. Light

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[1978Gui] Guinet, P., Halotier, D., Perroud, P., “Hydrogen Storage by Means of Reversible

Magnesium Alloys”, Eur. Communities [Rep] EUR1978, EUR 6085, Semin. Hydrogen

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62, 407-416 (1978) (Experimental, Magn. Prop., 36)

[1980Min] Mintz, M.H., Gavra, Z., Kimmel, G., Hadari, Z., “The Reaction of Hydrogen with

Magnesium Alloys and Magnesium Intermetallic Compounds”, J. Less-Common Met., 74,

263-270, (1980) (Experimental, 16)

[1980Sch1] Schefer, J., Fischer, P., Haelg, W., Stucki, F., Schlapbach, L., Didisheim, J.J., Yvon, K.,

Andresen, A.F., “New Structure Results for Hydrides and Deuterides of the Hydrogen

Storage Material Mg2Ni”, J. Less-Common Met., 74, 65-73 (1980) (Experimental, Crys.

Structure, 10)

[1980Sch2] Schefer, J., Fischer, P., Stucki, F., Schlapbach, L., Yvon, K., Didisheim, J.-J., “Neutron and

X-Ray Diffraction Investigations on Hydrided Mg2Ni” (in German), Helv. Phys. Acta, 53,

618 (1980) (Crys. Structure, Experimental, Abstract, 3)

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Transformation and Thermal Expansion of Mg-Ni alloys in a Hydrogen Atmosphere”,

J. Less-Common Met., 88, 57-61, (1982) (Crys. Structure, Experimental, 10)

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J. Less-Common Met., 88, 63-71 (1982) (Crys. Structure, Experimental, 10)

[1983Sem] Semenenko, K.N., Verbitskii, V.N., Kochukov, A.V., Sytnikov, A.N., “Interaction with

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Structure, 12)

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Three Polymorphs of Mg2NiH4”, J. Less-Common Met., 103, 277-283 (1984) (Crys.

Structure, 10)

[1984Nor] Noreus, D., Werner, P.E., “Structural Studies of Monoclinic Mg2NiD4”, J. Less-Common

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the System Mg2Ni-H2; Study by NMR of Mg2NiH0.3 and Mg2NiH4” (in French), J. Solid

State Chem., 52, 1-11 (1984) (Crys. Structure, Experimental, 8)

[1985Lue1] Luedecke, C.M., Deublein, G., Huggins, R.A., “Electrochemical Investigation of Hydrogen

Storage in Metal Hydrides”, J. Electrochem. Soc., 132(1), 52-56 (1985) (Equi. Diagram,

Experimental, 29)

[1985Lue2] Luedecke, C.M., Deublein, G., Huggins, R.A., “Investigation of Metal Hydrides with

Thermodynamic Calculations and Electrochemical Experiments”, Int. Symp. on Hydrogen

Systems, 1, 363-377 (1985) (Equi. Diagram, Experimental, 29)

322


MSIT®

H–Mg–Ni

[1985Nay] Nayeb-Hashemi, A.A., Clark, J.B., “The Mg-Ni (Magnesium-Nickel) System”, Bull. Alloy

Phase Diagrams, 6(3), 238-244 (1985) (Equi. Diagram, Review, #, 30)

[1985Nor] Noreus, D., Jansson, K., Nygren, M., “Structural, Thermal and Electrical Properties

Indicating Covalent Bonding in Mg2NiH4”, Z. Phys. Chem., 56, 191-202 (1985)


[1985Son] Song, M.Y., Pezat, M., Darriet, B., Hagenmuller, P., “A Kinetic Study on the Reaction of

Hydrogen with Mg2Ni”, J. Solid State Chem., 56, 191-202 (1985) (Experimental,

Thermodyn., 25)

[1986Nor] Noreus, D., Kihlborg, L., “Twinning at the Unit Cell Level in the Low Temperature Phase

of Mg2NiH4 Studied by Electron Microscopy”, J. Less-Common Met., 123, 233-239 (1986)


[1986Zol] Zolliker, P., Yvon, K., Baerlocher, Ch., “Low-Temperature Structure of Mg2NiH4:

Evidence for Microtwinning”, J. Less-Common Met., 115, 65-78 (1986) (Crys. Structure,

Experimental, 17)

[1987Lue] Luedecke, C.M., Deublein, G., Huggins, R.A., “Thermodynamic Characterization of Metal

Hydrogen Systems by Assessment of Phase Diagrams and Electrochemical Measurements”,

Int. J. Hydrogen Energy, 12(2), 81-88 (1987) (Thermodyn., Experimental, #, 18)

[1987Pos] Post, M.L., Murray, J.J., “Mg2Ni Hydride: in situ Heat Conduction Calorimetry of the Phase

Transition near 510 K”, J. Less-Common Met., 134, 15-26 (1987) (Thermodyn.,

Experimental, 35)

[1987San] San-Martin, A., Manchester, F.D., “The H-Mg (Hydrogen-Magnesium) System”, Bull.

Alloy Phase Diagrams, 8(5), 431-437 (1987) (Equi. Diagram, Review, #, 62)

[1987Sch] Schuermann, E., Sittard, M., Voelker, R., “Equivalent Influence of Alloying Elements on

the Nitrogen and Hydrogen Solubility in Liquid Ternary and Multicomponent Nickel-Base

Alloys” (in German), Z. Metallkd., 78(6), 457-466 (1987) (Experimental, 53)

[1987Sen] Senegas, J., Song, M.Y., Pezat, M., Darriet, B., “Phase Characterization and Hydrogen

Diffusion Study in the Mg-Ni-H system”, J. Less-Common Met., 129, 317-326 (1987)

(Experimental, 28)

[1988Sel] Selvam, P., Viswanathan, B., Swamy, C.S., Srinivasan, V., “Thermal Studies on the

Mg2NiH4: Existence of Additional Hydride Phase in the Mg2Ni-Hydrogen System”,

Thermochim. Acta, 125, 1-8 (1988) (Thermodyn., Experimental, 11)

[1989Hua1] Huang, N., Wang, Q.-D., Wu, J., Tang, J.-C., “Determination of Hydrogen Site Occupancy

in Cubic Mg2NiH4 by SCF-X -SW Calculations and Free Energy Considerations”, Z. Phys.

Chem., 163(1), 207-212 (1989) (Crys. Structure, Thermodyn., 12)

[1989Hua2] Hang, N., Yamauchi, H., Wu, J., Wang, Q.-D., “Determination of Hydrogen Atom

Configurations in Mg2NiH4 by Entropy Considerations and Infrared Spectra”, Z. Phys.

Chem., 163(1), 225-230 (1989) (Thermodyn., Experimental, 10)

[1989Way] Wayman, M.L., Weatherly, G.C., “The H-Ni (Hydrogen-Nickel) System”, Bull. Alloy

Phase Diagrams, 10(6), 569-580 (1989) (Equi Diagram, Review, #, 79)

[1993Jac] Jacobs, M.H.G., Spencer, P.J., “Thermodynamic Evaluation of the Systems Cu-Si, Mg-Ni,

Si-Sn and Si-Zn”, J. Chim. Phys., Physicochim. Biol., 90(2), 167-173 (1993)

(Thermodyn., 26)

[1994Bon] Bonhomme, F., Cerny, R., Fischer, P., Gingl, F., Hewat, A., Huang, B., Stetson, N.,

Yoshida, M., Yvon, K., Zolliker, M., “Powder Diffraction as a Routine Tool for ab-initio

Structure Determinations of Metals Hydrides”, Mater. Sci. Forum, 166-169, 597-602

(1994) (Crys. Structure, Review, 29)

[1996Ori] Orimo, S., Fujii, H., “Hydriding Properties of the Mg2Ni-H System Synthesized by

Reactive Mechanical Grinding”, J. Alloys Compd., 232, L16-L19 (1996) (Experimental, 14)

[1997Rud] Ruda, M., Abriata, J., “Thermodynamical Analysis of the Mg-Ni-H System” (in Spanish),

An. Assoc. Fis. Argent., 4 (1997) (Review, Thermodyn., 11)

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H–Mg–Ni

[1998Kom] Komaki, M., Amano, M., Numata, H., Nishimura, C., “Effects of Oxidation of Mg-Ni-H

and Mg-Ni-Al-H Alloys on Hydrogen Thermal Desorption Characteristics” (in Japanese),

J. Jpn. Inst. Met., 62(1), 111-116 (1998) (Experimental, 12)

[1998Li] Li, L., Akiyama, T., Kabutomori, T., Terao, K., Yagi, J., “In Situ X-ray Diffraction Study

of the Hydriding Combustion Synthesis of Mg2NiH4”, J. Alloys Compd., 281, 175-180


[1998Rud] Ruda, M., Abriata, J., “Thermodynamic Analysis and Amorphous Formation in the

Mg-Ni-H System”, XII WHEC, 10 pp., (1998) (Equi. Diagram, Thermodyn, #, 21)

[1998Son] Song, M.Y., Park, H.R., “Pressure-Composition Isotherms in the Mg2Ni-H2 System”,

J. Alloys Compd., 270, 164-167 (1998) (Experimental, 11)

[1999Deh] Dehouche, Z., Djaozandry, R., Goyette, J., Bose, T.K., “Evaluation Techniques of Cycling

Effect on Thermodynamic and Crystal Structure Properties of Mg2Ni Alloy”, J. Alloys

Compd., 288, 269-276 (1999) (Experimental, Crys. Structure, 14)

[1999Gar] Garcia, G.N., Abriata, J.P., Sofo, J.O., “Calculation of the Electronic and Structural

Properties of Cubic Mg2NiH4”, Phys. Rev. B, 59(18), 11746-11754 (1999) (Calculation,

Crys. Structure, Thermodyn., 34)

[1999Roe] Roennebro, E., Jensen, J.O., Noreus, D., Bjerrum N.J., “Structural Studies of Disordered

Mg2NiH4 Formed by Mechanical Grinding”, J. Alloys Compd., 293/295, 146-149 (1999)


[1999Zal] Zaluska, A., Zaluski, L., Stroem-Olsen, J.O., “Synergy of Hydrogen Sorption in Ball-Milled

Hydrides of Mg and Mg2Ni”, J. Alloys Compd., 289, 197-206 (1999) (Experimental, 22)

[1999Zen] Zeng, K., Klassen, T., Oelerich, W., Bormann, R., “Thermodynamic Analysis of the

Hydriding Process of Mg-Ni Alloys”, J. Alloys Compd., 283, 213-224 (1999) (Equi.

Diagram, Thermodyn., #, *, 41)

[2000Aiz] Aizawa, T., “Solid-State Synthesis of Magnesium Base Alloys”, Mater. Sci. Forum,

350/351, 299-310 (2000) (Review, 22)

[2000Kak] Kakuta, H., Kamegawa, A., Takamura, H., Okada, M., “Thermal Stability of Hydrides of

Magnesium-Transition Metal System Prepared under a High Pressure”, Mater. Sci. Forum,

350/351, 329-332 (2000) (Crys. Structure, Experimental, 11)

[2000Kuj] Kuji, T., Nakano, H., Aizawa, T., “Thermodynamic Properties of Mg2.0Ni Hydrides”,

Mater. Sci. Forum, 350/351, 311-314 (2000) (Experimental, Thermodyn., 9)

[2000Li] Li, L., Akiyama, T., Yagi, J., “Hydrogen Storage Alloy of Mg2NiH4 Hydride Produced by

Hydriding Combustion Synthesis from Powder of Mixture Metal”, J. Alloys Compd., 308,


[2000Wan] Wang, A.M., Ding, B.Z., Zhang, H.F., Hu, Z.Q., “Mechanical Alloying of Mg-33Ni Under

Hydrogen Atmosphere”, J. Mater. Sci. Lett., 19, 1089-1091 (2000) (Crys. Structure, 11)

[2001Kuj] Kuji, T., “Hydrogen Absorption Properties of Mg-Ni Alloys Prepared by Bulk Mechanical

Alloying”, Met. Mater. Int., 7(2), 169-173 (2001) (Experimental, 9)

[2001Li] Li, L., Aliyama, T., Yagi, J., “Activity and Capacity of Hydrogen Storage Alloy Mg2NiH4

Produced by Hydriding Combustion Synthesis”, J. Alloys Compd., 316, 118-123 (2001)


[2001Per] Perrot, P., Schmid-Fetzer, R., “H-Mg (Hydrogen-Magnesium)”, in “Ternary Alloys: A

Comprehensive Compendium of Evaluated Constitutional Data and Phase Diagrams”,

Vol. 18, Effenberg, G., Aldinger, F., Rogl, P. (Eds.), MSI GmbH, Stuttgart, 3-4 (2001) (Equi.

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Mg2NiH4”, J. Appl. Phys., 91(8), 5141-5148 (2002) (Crys. Structure, Electr. Prop.,

Experimental, 34)

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[2002Che] Chen, J., Yang, H.-B., Xia, Y.-Y., Kuriyama, N., Xu, Q., Sakai, T., “Hydriding and

Dehydriding Properties of Amorphous Magnesium - Nickel Films Prepared by a Sputtering

Method”, Chem. Mater., 14(7), 2834-2836 (2002) (Crys. Structure, Experimental,

Thermodyn., 21)

[2002Her] Herbst, J.F., “On Extending Miedema’s Model to Predict Hydrogen Content in Binary and

Ternary Hydrides”, J. Alloys Compd., 337, 99-107 (2002) (Calculation, Thermodyn., 20)

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Mg2-xNi (x=0-0.5) Alloys Prepared by Bulk Mechanical Alloying”, J. Alloy Compd.,

330/332, 590-596 (2002) (Crys. Structure, Experimental, #, 15)

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Products in Hydriding Combustion Synthesis of Mg2NiH4”, J. Alloys Compd., 345, 189-195

(2002) (Experimental, Kinetics, 16)

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Thermodynamic, Electronic, and Optical Properties of Monoclinic Mg2NiH4”, J. Appl.

Phys., 91(8), 4879-4885 (2002) (Crys. Structure, Optical Prop., Thermodyn., 31)

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Spectroscopy of Tetrahedral Ternary Metal Hydrides: Mg2NiH4, Rb3ZnH5 and their

Deuterides”, Phys. Chem. Chem. Phys., 4, 1732-1737 (2002) (Crys. Structure,

Experimental, 23)

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in a Mg-Ni-H System Prepared under an Ultra High Pressure”, J. Alloys Compd., 330/332,


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J. Alloys Compd., 356/357, 490-493 (2003) (Experimental, Phase Relation, Kinetics, 12)

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Mg2NiH4 under High Hydrogen Pressure”, J. Alloys Compd., 356/357, 697-700 (2003)


[2004Bob] Bobet, J.-L., Grigova, E., Khrussanova, M., Khristov, M., Stefanov, P., Peshev, P., Radev,

D., “Hydrogen Sorption Properties of Graphite Modified Magnesium Nanocomposites

Prepared by Ball Milling”, J. Alloys Compd., 366, 298-302 (2004) (Experimental,

Kinetics, 18)



Pearson Symbol/ Space Group/Prototype

Lattice Parameters[pm]

Comments/References

(Mg)< 650

hP2P63/mmcMg

a = 320.944c = 521.076

pure Mg at 25°C [V-C2][Mas2]

(Ni)< 1455

cF4Fm3mCu

a = 352.41 [Mas2, V-C2]

MgH2< 280

tP6P42/mnmTiO2 (Rutile)

a = 451.68c = 302.05

= 2.56 g cm-3 [1983Sem]

MgNi2< 1146

hP24P63/mmcMgNi2

a = 482.4c = 1582.6

[V-C2]

MgNi2 - - High pressure phase (5 MPa, 800°C) [2000Kak]

325


MSIT®

H–Mg–Ni

Mg2NiHx

Mg2Ni < 759

hP18P6222Mg2Ni

a = 523.15c = 1340.4a = 521.97c = 1324.33

at x = 0.3 [1982And]

[1982And]

NiH cF8 a = 372.51 0.08

a = 373.18

NiH0.9 in equilibrium with Ni [1989Way]NiH [1989Way]

* Mg2NiH4(h2)> 245

cF28Fm3mCaF2

a = 648.9 = 2.61 g cm-3

[1980Sch1, 1982Ono2]

* Mg2NiH4(h1)245 - 235

oP* a = 1136b = 1116c = 912

= 2.56 g cm-3

[1980Min, 1984Dar]probably metastable

* Mg2NiH4(r)< 235

mC56C2/cK2PtCl6

a = 1299b = 639c = 659.8

= 93.22°

[1984Dar, 1994Bon]

* Mg2Ni3H3.4 oP* a = 885.9 0.4b = 1374.0 0.5c = 469.4 2

High pressure phase, 800°C, 5 GPa [2002Tak]


Pearson Symbol/ Space Group/Prototype

Lattice Parameters[pm]

Comments/References

350

100

150

200

250

300

Mg Ni2

1.0 MPa

18 kPa

50 kPa

0.13 MPa

0.33 MPa

Mg NiH (h )2 4 2

Mg NiH (r)2 4

Mg

Ni

H

66.67

33.33

00.00

Mg

Ni

H

28.57

14.29

57.14

Te

mp

era

ture

,°C

4020

H, at.%

Fig. 1: H-Mg-Ni.

Vertical section

Mg2Ni - Mg2NiH4

showing some

isobaric curves

326


MSIT®

H–Mg–Ni

1.0 MPa

0.1 MPa

0.01 MPa

100

80

60

40

20

10 20 30 40 50

89.0

28.1

8.90

34.8

11.0

3.48

Mg

xNi-3

·10

xH

-4·1

0

Fig. 2: H-Mg-Ni.


700°C of the Mg rich

corner under

hydrogen pressures of

0.01, 0.1 and 1MPa

20

40

60

80

20 40 60 80

20

40

60

80

Ni Mg

H Data / Grid: at.%

Axes: at.%

MgH2

MgNi2

Mg2Ni

Mg2NiH

4

Fig. 3: H-Mg-Ni.


section at 301°C and

3.47 bar [1999Zen]

327


MSIT®

Li–Mg–Si

Lithium – Magnesium – Silicon

Oksana Bodak and Rainer Schmid-Fetzer

Literature Data

The ternary Li-Mg-Si phase relations were established using high purity samples prepared by levitation

melting [2004Kev]. Starting materials were magnesium chips (99.98 mass%), lithium bulk material (99.9

mass%) and silicon chips (99.9998 mass%). The elements were weighed and mixed in a glove-box with Ar

atmosphere and pressed under a pressure of 100 MPa into small pellets of around 0.5 g. Alloys were

prepared from that by levitation melting under purified argon atmosphere. Heating power was controlled

carefully to avoid evaporation. Weight loss was found less than 1 mass% after levitation melting. Analysis

was done using X-ray powder diffraction (XRD), optical metallography and differential thermal analysis

(DTA) in customized sealed tantalum crucibles. The obtained X-ray diffraction patterns were analyzed

quantitatively in comparison with simulated X-ray spectra. For DTA analysis the samples were sealed under

pure argon at 1 bar in specially adapted tantalum containers using electric arc welding. Evaporation and

oxidation of the samples was completely avoided with this technique. This procedure enables reproducible

DTA signals that cannot be obtained otherwise. After thermal analysis, the alloys were again examined by

XRD.

A thermodynamic assessment of the ternary Li-Mg-Si system was worked out in parallel to the experimental

study and used to select the relatively small number of eight key alloys [2004Kev]. The experimentally well

supported calculated phase equilibria in the entire composition and temperature range of the Li-Mg-Si

system are presented, including the liquidus surface and invariant reactions [2004Kev].

The few earlier experimental investigations [1986Nes, 1992Pav2, 1992Pav3] presented altogether five

ternary phases of various compositions. The isothermal section of the phase diagram at 200°C was studied

by [1992Pav3, 1996Dmy]. They prepared alloys by arc melting in a purified Ar atmosphere using a Li-Mg

master alloy [1992Pav3]. Elemental metals were also used for the preparation of alloys with purity Li (98.2

mass%), Mg (99.98 mass%) and Si (99.9999 mass%). Annealing was carried out in Ta containers at 200°C

for 240 h. X-ray powder diffraction was used for the phase analysis. Four ternary phases Li5MgSi4,

Li12Mg3Si4, Li2MgSi and LiMg2Si were reported in this system [1992Pav3]. All the phases have no

significant solubility. In addition, no ternary solubility of binary phases was reported [1992Pav3], however,

[2004Kev] later showed that LiMg2Si and Mg2Si form a continuous solution. [1986Nes] reported the

monoclinic phase Li8MgSi6.

A study of alloys based on the Li-Mg system has been made concerning development of microstructure,

mechanical properties and oxidation behavior of Li-Mg-Si alloys [1990Sid]. Electrochemical Li insertion

into the anti-fluorite type Mg2Si was performed and the structural variation during the insertion was

examined [2000Mor].

Binary Systems

The binary system Li-Mg from [Mas2] is accepted. A recent thermodynamic assessment of the Mg-Si phase

diagram [2004Kev] given in Fig. 1 is accepted, that resolves unrealistic inverted liquid miscibility gaps in

previous thermodynamic descriptions. The binary system Li-Si from [Mas2] is modified in two

thermodynamic assessments given in [1995Bra]. In the first assessment of [1995Bra] the chemical potential

values for the solid two-phase regions are in good agreement with the experimental values. However, the

liquidus in the region of the Li12Si7 and Li7Si3 phases shows significant deviation from the experimental

points. In the second assessment [1995Bra], a better agreement with the experimental phase diagram data

is shown, but unrealistic values for the chemical potential values and thermodynamic parameters are

produced. In the present report, the first assessment of [1995Bra] is preferred because of the more realistic

Gibbs energies. The phase diagram calculated by [2004Kev] with these parameters is given in Fig. 2. The

composition of the Li-Si binary phases in [Mas2] differ slightly from those assigned for the phases in earlier

works. The phases named Li10Si3 and Li4Si possibly are the Li22Si5 phase for which two alternative

328


MSIT®

Li–Mg–Si

structures were proposed. The Li2Si phase, which is related closely in crystallographic structure to Li7Si3(or “Li14Si6”), is absent in the phase diagram of [Mas2]. However, in the investigation of the ternary system

both phases, Li2Si and Li7Si3, were observed [1992Pav3]. Crystal structure data of binary phases are also

given in Table 1 and former phase designations are provided.

Solid Phases

Crystal structure data on all solid phases are given in Table 1. Three ternary intermetallic phases were

confirmed in this system: 1 (Li8MgSi6), 2 (Li12Mg3Si4), and 3 (Li2MgSi). The LiMg2Si phase reported

previously is not a separate phase but the interstitial solid solution of Li in the binary Mg2Si phase, denoted

LixMg2Si, with a maximum of x 0.91 0.05 [2004Kev]. This was also corroborated by DTA

measurements. Also confirmed [2004Kev] was the previous finding [1992Pav2] that all ternary phases 1,

2, and 3, show negligible ranges of homogeneity. The insertion mechanism of Li in LixMg2Si was studied

by [2000Mor].

The monoclinic phase reported as Li5MgSi4 (Li50Mg10Si40) [1992Pav1] has a similar composition to

Li8MgSi6 (Li53.33Mg6.77Si40) [1986Nes, 1996Dmy]. It may be assumed that both compositions describe

the same phase and Li8MgSi6 is the real composition of the phase.

[1968Pau] failed to detect an fcc structure at the composition Li2MgSi, in contrast to the confirmed finding

of 3 (Li2MgSi). It is also stated in [1996Wen] that the perfect stoichiometry Li2MgSi does not exist,

without providing more details. However, a Li2-2xMg1-xSi (x~0.06) phase with a defect Li3Bi type structure

(space group Fm3m) is mentioned and was studied by theoretical calculations. Experimentally, however, in

the same paper Li2-2xMg1-xSi (x = 0.05) was reported to be two-phase and melting at 1017°C [1996Wen].


The invariant equilibria calculated from the experimentally supported thermodynamic description

[2004Kev] are given in Table 2.

Liquidus, Solidus and Solvus Surfaces

The calculated liquidus surface of the Li-Mg-Si phase diagram [2004Kev] is presented in Fig. 3. The

primary liquidus field of Li12Si7 is extremely small and located near the binary edge of the phase diagram.

Visible on the liquidus surface in that region is merely the primary field of 1.

Along the line e1-U7 it is shown [2004Kev] that the gradual transition from eutectic, L LixMg2Si + 3, to

peritectic, L+ 3 LixMg2Si, monovariant reaction type is intricate. It depends on the overall alloy

composition in the three-phase field and may occur at different temperatures. The classical tangent criterion

fails, since not only the liquid but also the LixMg2Si solid composition moves with temperature.

Isothermal Sections

Calculated isothermal sections at 600, 400 and 200°C are given in Figs. 4 - 6. The solid state equilibria at

200°C in Fig. 6 are partially also supported by the findings of [1992Pav3], noting that instead of

“LiMg2Si+Mg2Si” only the single phase range LixMg2Si exists. Another difference is that 3 coexists with

Li-Si phases whereas a conflicting tie line 2+ 1 was reported by [1992Pav2]. However, the X-ray spectra

of these phase assemblies are very complex and interpretation may be inconclusive. In addition, the

thermodynamic modeling [2004Kev] had shown that a 2+ 1 tie line could only be modeled with a much

lower stability of 3, resulting in a loss of the key equilibrium 3+(Mg). This 3+(Mg) equilibrium is firmly

established by the XRD analysis of a slowly cooled sample, the microstructure of the as-cast sample, and

the secondary DTA effect at U7.


A calculated partial vertical section at 10 at.% Si is shown in Fig. 7. Three additional vertical sections are

given in [2004Kev].

329


MSIT®

Li–Mg–Si

Thermodynamics

A consistent Calphad-type thermodynamic analysis of the Li-Mg-Si system is given by Kevorkov et al.

[2004Kev]. It is well supported by the experimental results obtained at 200°C and the phase assemblies and

transition temperatures observed in slowly cooled samples and phase assemblies of as-cast samples. No

ternary parameters were used for the liquid phase.


Mg2Si is currently under consideration as anodic material for Li-ion-cells. The Li-Mg-Si phase diagram is

important for an understanding of the electrochemical reactions of Li-incorporation during charging.

It was found that lithium insertion into Mg2Si proceeds stepwise according to the following reactions:

Mg2Si+2Li++2e- Li2MgSi+Mg, and Mg + yLi++ye- LiyMg [2000Mor].

Solidification and phase transformations occurring in samples slightly off the stoichiometry of the ternary

intermetallic phases are rather intricate. It is demonstrated that this intricate behavior can be quantitatively

understood using a number of calculated ternary phase diagram sections and invariant reactions [2004Kev].

This knowledge may be also important in the processing of these intermetallic phases and the control of

trace amounts of foreign phases.

If lithium is added to magnesium alloys containing small amounts of Si, it dissolves at a higher fraction in

the Mg2Si phase compared to the (Mg) matrix. At higher lithium-addition, the ternary phase 3 (Li2MgSi)

starts forming.

Miscellaneous

Wengert et al. [1996Wen] present Car-Parrinello molecular dynamics simulations of a novel superionic

conductor, Li2-2xMg1-xSi (x ~0.06), at different temperatures. The calculations clarify the nature of the ionic

conduction and lead to the prediction of the first inorganic magnesium superionic conductor. Both lithium

and magnesium are found to act as charge carries [1996Wen].

References

[1968Pau] Pauly, H., Weiss, A., Witte, H., “Face Centred Cubic Alloys of Composition Li2MgX with

Body-Centred Substructure” (in German), Z. Metallkd., 59(5), 414-418 (1968) (Crys.


[1986Nes] Nesper, R., Curda, J., Schnering, H.G., “Li8MgSi6 a Novel Zintl Compound Containing

Quasi-Atomic Si5 Rings”, J. Solid State Chem., 62, 199-206 (1986) (Crys. Structure,

Experimental, 20)

[1990Sid] Siddhartha, D., “A Study of Alloys Based on the Mg-Li System”, Diss. Abstr. Int. B, 50(7),

3120-B (1990) (Equi. Diagram, Experimental, Mechan. Prop., 0)

[1992Pav1] Pavlyuk, V.V., Dmytriv, G.S., Starodub, P.K., “Crystal Structure of the Compounds of the

Li-M-X (M = Mg, Al; X = Si, Ge, Sn) Systems” (in Russian), Cryst. Chem. Inorg. Coord.

Compounds, VI Conf. September 1992, L’viv (Abstract), 210 (1992) (Crys. Structure,

Experimental, 0)

[1992Pav2] Pavlyuk, V.V., Bodak, O.I., “Crystal Structure Of Lithium Magnesium Silicide

(L12Mg3Si4) and Lithium Aluminum Silicide (Li12Al3Si4)” (in Russian), Neorgan. Mater.,

28(5), 988-990 (1992) (Crys. Structure, Experimental, 3)

[1992Pav3] Pavlyuk, V.V., Bodak, O.I., Dmytriv, G.S., “Interaction of components in the Li-

(Mg,Al)-Si Systems” (in Russian), Ukr. Khim. Zh., 58(9), 735-737 (1992) (Equi. Diagram,

Crys. Structure, Experimental)

[1995Bra] Braga, M.H., Malheiros, L.F., Ansara, I., “Thermodynamic Assessment of the Li-Si

System”, J. Phase Equilib., 16(4), 324-330 (1995) (Equi. Diagram, Thermodyn.,

Assessment, Calculation, #, 15).

330


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Li–Mg–Si

[1996Dmy] Dmytriv, G.S., “Phase Equilibria and Crystal Structure of Compounds in the Systems

Mg-Li-Si, Ca-Li-{Si,Ge}, Al-Li-{Si,Ge,Sn}, Zn-Li{Al,Sn}” (in Ukrainian), Doctor Thesis,

Lviv State University, pp.22. (1996) (Equi. Diagram, Crys. Structure, Experimental, 10)

[1996Wen] Wengert, S., Nesper, R., Andreoni, W., Parrinello, M., “Ionic Diffraction in a Ternary

Superionic Conductor: an ab initio Molecular Dynamic Study”, Phys. Rev. Lett., 77(25),


[2000Mor] Moriga, T., Watanabe, K., Tsuji, D., Massaki, S., Nakabayashi, I., “Reaction Mechanism of

Metal Silicide Mg2Si for Li Insertion”, J. Solid State Chem., 153, 386-390 (2000) (Crys.


[2004Kev] Kevorkov, D., Schmid-Fetzer, R., Zhang, F., “Phase Equilibria and Thermodynamics of the

Mg-Si-Li System and Remodeling of the Mg-Si System”, J. Phase Equilib., 25(2), 140-151

(2004) (Equi. Diagram, Crys. Structure, Thermodyn., Assessment, Calculation,

Experiment, #, *, 23)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Li)

180.6 - (-193)

cI2

Im3m

W

a = 351.0

a = 351.4 to 349.5

pure Li at 25°C [V-C2]

dissolves 75.5 at.% Mg at 588°C [Mas2]

at 70 at.% Mg exists up to 592°C

30 to 70 at.% Li [V-C2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.944

c = 521.076

a = 320.95 to

319.18

c = 521.06 to

513.19

pure Mg at 25°C [V-C2]

dissolves 17 at.% Li at 588°C [Mas2]

0 to 18.4 at.% Li [V-C2]

(Si)

< 1414

cF8

Fm3m

C (diamond)

a = 543.09 pure Si at 25°C [V-C2]

Li22Si5< 619

cF432

F23

Li22Pb5

a = 1875 [V-C2] earlier “Li4Si”

[Mas2]. Melting [1995Bra], at 628°C

[Mas2]

Li13Si4< 732

oP34

Pbam

Li13Si4

a = 799

b = 1521

c = 443

[V-C2] earlier “Li7Si2”

[Mas2]. Melting [1995Bra], at 722°C

[Mas2]

Li7Si3< 748

hR21

R3m

Mo2B5

a = 443.5

c = 1813.4

[V-C2] denoted as “Li14Si6”

[1992Pav3]. Melting [1995Bra], at

752°C [Mas2]

“Li2Si” mC12

C2/m

OsGe2

a = 770

b = 441

c = 656

= 113.4°

[V-C2]. Not considered as stable binary

phase.

Li12Si7< 629

oP152

Pnma

Li12Si7

a = 861.0

b = 1973.8

c = 1434.1

[V-C2] earlier “Li13Si7”.

Melting [1995Bra], at 648°C [Mas2]

331


MSIT®

Li–Mg–Si


LixMg2Si

Mg2Si

< 1081

cF16 (cF12)

F43m (Fm3m)

MnCu2Al (CaF2)

a = 635 to 638

a = 634.7 0.4

a = 638.8

x = 0 to 0.91 [2004Kev]

x = 0, [V-C2], Melting [2004Kev]

“LiMg2Si”[1992Pav3, 1992Pav1,

1996Dmy]

* 1, Li8MgSi6< 683

mP48

P21/m

Li8MgSi6

a = 1270.1

b = 434.7

c = 1050.7

= 107.58°

[1986Nes, 2004Kev]

Denoted “Li5MgSi4” [1992Pav1,

1992Pav3]

* 2, Li12Mg3Si4< 674

cI76

I43d

Li12Mg3Si4

a = 1068.8 [1992Pav2, 1992Pav3, 1996Dmy,

2004Kev]

* 3, Li2MgSi

< 995

cF16

F43m

CuHg2Ti

a = 637.0 [1992Pav1, 1992Pav3, 1996Dmy,

2004Kev]


Li Mg Si

L 3 995 congruent L 50 25 25

L 3+ LixMg2Si 968 e1 L 37.26 36.91 25.84

L 3+ Li7Si3 741 e2 L 68.85 1.31 29.84

L 3+ Li13Si4 721 e3 L 74.59 1.91 23.51

L 3+ Li13Si4 + Li7Si3 719 E1 L 73.37 1.58 25.04

L + LixMg2Si (Si) + 3 696 U1 L 49.76 4.96 45.28

L + 3 1 683 p1 L 54.06 2.04 43.89

L + 3 1 + (Si) 679 U2 L 51.58 3.45 44.97

L + 3 2 675 p2 L 79.45 4.40 16.15

L + 3 2 + Li13Si4 675 U3 L 79.47 4.27 16.26

L + 3 Li7Si3 + 1 660 U4 L 60.64 0.76 38.60

L + Li7Si3 1 + Li12Si7 629 U5 L 59.52 0.11 40.36

3 + (Si) 1 + LixMg2Si 608 U6 - - -

L 1 + Li12Si7 + (Si) 604 E2 L 56.68 0.03 43.29

L + LixMg2Si 3 + (Mg) 597 U7 L 18.77 80.05 1.18

L + Li13Si4 2 + Li22Si5 592 U8 L 83.84 5.05 11.11

L 3 + (Li) 584 e4 L 27.67 71.19 1.13

L 3 + (Mg) + (Li) 583 E3 L 23.79 75.29 0.91

L + 3 2 + (Li) 455 U9 L 62.74 36.12 1.13

L + 2 Li22Si5 + (Li) 318 U10 L 80.27 19.21 0.51

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

332


MSIT®

Li–Mg–Si

20 40 60 80

0

250

500

750

1000

1250

1500

Mg Si

Si, at.%

Te

mp

era

ture

, °C

L

(Si)

(Mg)

Mg2Si

L+(Si)

L+Mg2Si

946.5°C

637.4°C

54.12

1.45

1081.4°C

(Mg)+Mg2Si

Mg2Si+(Si)

Fig. 1: Li-Mg-Si.

Calculated Mg-Si

phase diagram

20 40 60 80

0

250

500

750

1000

1250

1500

Li Si

Si, at.%

Te

mp

era

ture

, °C

180.3°C

604°C619°C

732°C 748°C

629°C

(Li)

(Si)

43.28

Li22Si5

Li13Si4

Li7Si3Li12Si7

L

Fig. 2: Li-Mg-Si.

Calculated Li-Si

phase diagram

333


MSIT®

Li–Mg–Si

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%

LixMg

2Si

U9

U7

E3

e4Li

22Si

5U

10

(Li) (Mg)

e1

U2

Li7Si

3 p1

U4

e3

Li13

S4

U3

p2

Li12

Si7

τ1

τ2

τ3

(Si)

U1

U8

E1

e2

τ1

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%(Si)

τ1

τ3

τ2

L

Mg2Si

Li12

Si7

Li7Si

3

Li13

Si4

Li22

Si5

(Mg)

τ1+Li

xMg

2Si+(Si)

τ1+τ

3+Li

xMg

2Si

τ3+L+Li

xMg

2Si

Fig. 3: Li-Mg-Si.

Calculated liquidus

surface

Fig. 4: Li-Mg-Si.


section at 600°C

334


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Li–Mg–Si

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%

τ1

τ3

τ2

Mg2Si

(Mg)L (Li)

Li12

Si7

Li7Si

3

Li13

Si4

Li22

Si5

(Si)

τ1+Li

xMg

2Si+(Si)

τ1+τ

3+Li

xMg

2Si

τ3+(Mg)+Li

x Mg2Si

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%(Si)

Mg2Si

(Li) (Mg)

τ1

τ3

τ2

Li12

Si7

Li7Si

3

Li13

Si4

Li22

Si5

L

τ1+Li

xMg

2Si+(Si)

τ1+τ

3+Li

xMg

2Si

τ3+(Mg)+Li

xMg2Si

Fig. 6: Li-Mg-Si.


section at 200°C

Fig. 5: Li-Mg-Si.


section at 400°C

335


MSIT®

Li–Mg–Si

10 20

250

500

750

1000

Li 0.00

Mg 90.00

Si 10.00

Li 30.00

Mg 60.00

Si 10.00Li, at.%

Te

mp

era

ture

, °C

L+LixMg2Si

L+LixMg2Si+Li2MgSi

(Mg)+LixMg2Si+Li2MgSi

(Mg)+LixMg2Si

(Mg)+Li2MgSi

U7

LFig. 7: Li-Mg-Si.

Calculated partial

vertical section at

constant 10 at.% Si

336


MSIT®

Li–Mg–Zn

Lithium – Magnesium – Zinc

Volodymyr Pavlyuk, Oksana Bodak, Hans Leo Lukas

Literature Data

Early works [1950Bus, 1956Jon] in this system concern the effect of alloying elements on the mechanical

properties of the bcc (Li) solid solution phase of the Li-Mg binary system. The ductile binary phase is

considerably hardened by precipitations in Zn containing alloys. The precipitates were designated as

LiMgZn and Li2MgZn. The hardening is stable only below room temperature [1956Jon]. Weinberg et al.

[1956Wei] first investigated the ternary phase diagram, giving isothermal sections at 400, 300, 200, and

100°C excluding the ranges with less than 15 at.% Li or more than 60 at.% Zn. A ternary cubic phase, called

-LiMgZn, with a = 745 pm was found, whereas Li2MgZn was stated to be not an equilibrium phase. A

four-phase reaction L+LiMgZn (Li)+LiZn was found at 317°C. LiMgZn, having some homogeneity range

around a composition MgLi0.77Zn1.23, was later [1958Kry] identified as cubic C15 type Laves phase.

Alloys were typically prepared from Li (98.5 to 99.5 mass% purity), Mg (94.94 to 99.99 mass%) and Zn

(99.97 to 99.995 mass%) by melting in molybdenum or corundum crucibles under a LiCl-LiF flux.

[1958Sha] measured the simultaneous solubility of Li and Zn in the hcp (Mg) solid solution and [1960Sha]

that of Zn in the bcc (Li) solid solution with 67 to 71 at.% Mg. The results agree with [1956Wei] within the

scatter.

Pauly et al. [1968Pau] in a systematic study of Li2MgX alloys with 16 different elements X found Li2MgZn

to be not single phase, but they identified a NaTl type phase with a = 628 pm, Li in the Na positions and

Mg+Zn statistically distributed in the Tl positions of this structure type.

Between MgZn2 (C14 type) and MgLi0.77Zn1.23 (MgCu2-C15 type [1958Kry]) several other Laves phases

were detected and their structures determined [1958Kry, 1971Kry, 1971Mel, 1974Fai, 1974Kry1,

1974Kry2, 1974Mel1, 1974Mel2, 1974Mel3, 1975Mel].

[1974Yar] determined the crystal structure of Mg2Zn3 and found a ternary stacking variant of this structure

with composition Li25Mg24Zn.

An assessment of the ternary Li-Mg-Zn system has been reported by [1987Mal].

Binary Systems

The binary Li-Mg system is accepted from [1984Nay]. It shows no intermediate compounds but only the

(Mg based) and the very large (Li based) solid solutions. In the Mg-Zn system, accepted from [2001Shc],

there are five intermediate phases, Mg51Zn20 (high temperature phase), MgZn, Mg2Zn3, MgZn2 and

Mg2Zn11. The Li-Zn system is accepted from [1991Pel]; three high temperature intermediate phases have

been reported, LiZn4, Li2Zn5 and Li2Zn3 with ordering at lower temperatures. Li2Zn3 forms LiZn and

Li2Zn3 below about 175°C, Li2Zn5 transforms peritectoidally to Li2Zn5 and LiZn4 peritectoidally to

LiZn4. A room temperature phase, LiZn2 is formed peritectoidally at 93°C. The crystal structure is

reported only for LiZn.

Solid Phases

The Li-Mg-Zn system is characterized by the formation of a high number of Laves phases (Laves polytypes)

along the section at 33.3 at.% Mg. Their crystal data are reported in Table 1 following the summary of

[1974Mel2].

A metastable ternary phase, Li2MgZn is formed during aging of Zn containing (Li) solid solutions

[1950Bus, 1956Jon, 1956Wei, 1968Pau]. It has the same crystal structure (NaTl type) as LiZn, but is

distinguished by a different lattice parameter (see Table 1).

337


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Li–Mg–Zn


Only one invariant four-phase equilibrium is reported [1956Wei]:

L + LiMgZn (Li) + LiZn, 317°C.

Isothermal Sections

The isothermal sections at 400 and 300°C are shown in Figs. 1 and 2, respectively. Figure 1 is taken from

[1956Wei], some corrections are made to satisfy Schreinemaker’s rule. Figure 2 taken from [1974Kry2]

was completed by the Laves phase L3’, given in the summary of [1974Mel1]. It agrees fairly well with the

300°C isothermal section of [1956Wei]. LiZn was replaced by Li2Zn3 in order to agree with the accepted

binary Li-Zn system.

References

[1933Tar] Tarschisen, L., “X-ray Investigation of the MgZn and MgZn5 Compounds”, Z. Kristallogr.,

A186, 423-438 (1933) (Crys. Structure, 6)

[1950Bus] Busk, R.S., Leman, D.L., Casey, J.J., “The Properties of Some Magnesium - Lithium Alloys

Containing Zinc and Aluminum”, J. Metal. Trans. AIME, 188, 945-951 (1950)

(Experimental, 6)

[1956Jon] Jones, W.R.D., Hogg, G.V., “The Stability of Mechanical Properties of Beta-Phase

Magnesium - Lithium Alloys”, J. Inst. Met., 85, 255-261 (1956-57) (Experimental, 7)

[1956Wei] Weinberg, A.F., Levinson, D.W., Rostoker, W., “Phase Relations in the Mg-Li-Zn Alloys”,

Trans. Amer. Soc., 48, 855-871 (1956) (Equi. Diagram, Experimental, 6)

[1958Kry] Krypyakevich, P.I., “Crystal Structure of the Compounds MgLiZn and MgLi0.25Zn1.75” (in

Ukrainian), Khim. Zbirn. L'viv Univer., 5, 107-114 (1958) (Crys. Structure,

Experimental, 7)

[1958Sha] Shamray, F.I., Krylova, E.Ya., “On Reciprocal Solubility of Zn and Li in Mg in Solid State

at Different Temperatures” (in Russian), Tr. Inst. Met., Akad. Nauk SSSR, 2, 231-237 (1958)

(Equi. Diagram, Experimental)

[1959Ray] Raynor, G.V., “Intermediate Phases in Magnesium Alloys”, in “The Physical Metallurgy of

Magnesium and Its Alloys”, Pergamon Press, London, New York, Paris, Los Angeles,

145-215 (1959) (Review, Crys. Structure, 35)

[1960Sha] Shamray, F.I., Krylova, E.Ya., “On Reciprocal Solubility of Zn and Li in the Mg-Zn-Li

System -Phase in Solid State” (in Russian), Tr. Inst. Met., Akad. Nauk SSSR, 4, 200-207

(1960) (Equi. Diagram, Experimental)

[1968Pau] Pauly, H., Weiss, A., Witte, H., “Fcc Allyos of Composition Li2MgX with Body-Centred

Substructure” (in German), Z. Metallkd., 59(5), 414-418 (1968) (Crys. Structure,

Experimental, 15)

[1971Kry] Krypyakevich, P.I., Melnik, E.V., “Nine-Layer Laves's Phases in Mg-Li-Zn, Mg-Cu-Zn and

Mg-Co-Ni” (in Ukrainian), Dop. Akad. Nauk Ukr. RSR, Ser. A, 11, 1046-1048 (1971) (Crys.


[1971Mel] Melnik, E.V., Krypyakevich, P.I., “New Structural Types of Laves's Phases and Their

Representatives” (in Russian), Vsesoyuzn. Konf. Kristallokhim. Intermetal. Soed., L'viv, 27

(1971) (Crys. Structure, Experimental)

[1971Tes] Teslyuk, M.Yu., Mitrofanova, M.F., Melnik, E.V., Malinkovich, A.N., “Study of Structural

Transformation with Aging of Two-Phase ( + ) Alloys of the Mg-Li-Zn System” (in

Russian), in “Struktura i Svoistva Legkikh Splavov”, 40-44 (1971) (Experimental, 9)

[1974Fai] Fainshtein, G.S., “Morphotropy of Laves Phase Transformation l1(2H)-l3(4H) in Ternary

Systems of Non-Transition Metals” (in Russian), Metallofizika, 52, 80-84 (1974) (Crys.


338


MSIT®

Li–Mg–Zn

[1974Kry1] Krypyakevich, P.I., Melnik, E.V., “Structure of 14-Layers Laves's Phase (MgLi0.11Zn1.89)”

(in Ukrainian), Dop. Akad. Nauk Ukr. RSR, Ser. A, 9, 847-849 (1974) (Crys. Structure,

Experimental, 15)

[1974Kry2] Krypyakevich, P.I., Melnik, E.V., “New Results on Crystal Chemistry of Multilayer Laves's

Phases” (in Russian), Metallofizika, 52, 71-74 (1974) (Equi. Diagram, Crys. Structure,


[1974Mel1] Melnik, E.V., “Laves's Phases in Mg-Li-Zn” (in Ukrainian), Dop. Akad. Nauk Ukr. RSR,

Ser. A, 10, 949-953 (1974) (Crys. Structure, Experimental, 18)

[1974Mel2] Melnik, E.V., Krypyakevich, P.I, “Mg2LiZn3 Compound, Laves's Phase with a New Type

of Superstructure” (in Russian), Kristallografiya, 19(3), 645-646 (1974) (Crys. Structure,

Experimental, 8)

[1974Mel3] Melnik, E.V., Kripyakevich, P.I., “Concentration - Dependent Polytypes and Structure

Formation in Mg(Li,Zn)2” (in Russian), Tezisy Dokl. Vses. Konf. Kristallokhim. Intermet.

Soeden, 2nd, Lvov, 17-18 (1974) (Crys. Structure, 1)

[1974Yar] Yarmolyuk, Ya.P., Kripyakevich, P.I., Melnik, E.V.,“New Hybrides of Zr4Al3 and Laves

Phase Structure Types, Mg4Zn7, Mg7(Li,Zn)9)” (in Russian), Tezisy Dokl. Vses. Konf.

Kristallokhim. Intermet. Soeden, 2nd, Lvov, p.16 (1974) (Crys. Structure, 0)

[1975Mel] Melnik, E.V., “Investigations of the Ternary Systems Mg-Li-{Zn, Cd, Sn, Cu, Mn, La, Ce}

and Problems of Crystal Chemistry of Formed Compounds” (in Russian), Summery Diss.

Kand. Khim. Nauk, L'vov, 15 (1975) (Equi. Diagram, Crys. Structure, Review, #)

[1975Yar] Yarmolyuk, Ya.P., Kripiakevich, P.I., Melnik, E.V., “The Crystal Structure of the Mg4Zn7

Compound”, Sov. Crystallogr., 20(3), 538-542 (1975), translated from Kristallografiya,

20(3), 538-542 (1975), (Experimental, Crys. Structure, 16)

[1981Hig] Higashi, I., Shiotani, N., Uda, M., Mizoguchi, T., Katoh, H., “The Crystal Structure of

Mg51Zn20”, J. Solid State Chem., 36, 225-233 (1981) (Crys. Structure, Experimental, 11)

[1984Nay] Nayeb-Hashemi, A.A., Pelton, A.D., Clark, G.B., “The Li-Mg (Lithium - Magnesium)

System”, Bull. Alloy Phase Diagrams, 5(4), 365-374 (1984) (Equi. Diagram, Review, #, 9)

[1987Mal] Mallik, A.K., “The Lithium - Magnesium - Zinc System”, J. Alloy Phase Diagrams, 3,

12-15 (1987) (Equi. Diagram, Crys. Structure, Review, 9)

[1991Pel] Pelton, A.D., “The Li-Zn (Lithium - Zinc) System”, J. Phase Equilib., 12(1), 42-45 (1991)

(Equi. Diagram, Review, #, 12)




[1994Goe] Goedecke, T., Sommer, F., “Stable and Metastable Phase Equilibria in MgZn2-Zn and

Mg2Sn-MgZn2-Sn-Zn Alloys” (in German), Z. Metallkd., 85(10), 683-691 (1994)

(Experimental, Equi. Diagram, #, 9)

[2001Shc] Shcherban, O., Ilyenko, S., “Mg-Zn (Magnesium-Zinc)” in “Ternary Alloys: A

Comprehendium of Evaluated Consitutional Data and Phase Diagrams”, Effenberg, G.,

Aldinger, F., Rogl, P. (Eds.), MSI GmbH, Stuttgart, (2001) (Crys. Structure, Equi. Diagram,

Assessment, 9)

339


MSIT®

Li–Mg–Zn


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Li)

< 180.6

cI2

Im3m

W

a = 351.0

a = 351.4 to 349.5

pure Li at 25°C [V-C2]


at 70 at.% Mg exists up to 592°C

dissolves 1.5 at.% Zn at 161°C [Mas2]

30 to 70 at.% Li [V-C2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.944

c = 521.076

a = 320.95 to

319.18

c = 521.06 to

513.19

a = 320.99 to

319.57

c = 521.08 to

518.82



at 0 to 18.4 at.% Li [V-C2]


at 0 to 2.81 at.% Zn [V-C2]

(Zn)

< 419.5

hP2

P63/mmc

Mg

a = 266.47

c = 494.69

a = 266.55 to

266.55

c = 494.88 to

494.67

pure Zn at 25°C [V-C2]



at 0 to 1 at.% Li [V-C2]

Li2Zn3(h)

520 - 155

- - ~50-67 at.% Zn [Mas2]

LiZn

< 177

cF16

Fd3m

NaTl

a = 622 [1991Pel]

Li2Zn3(r)

< 174

c* ? a = 427 [1991Pel]

Li2Zn5(h)

502 - 168

- - [1991Pel]

Li2Zn5(r)

< 268

h* ? - [1991Pel]

LiZn2

< 93

- - [1991Pel]

LiZn4(h)

461 - 65

hP2 a = 278

c = 439

[1991Pel]

LiZn4(r)

< 245

h* ordered [1991Pel]

Mg51Zn20

341.1 - 325

oI158

Immm

Mg51Zn20

a = 1408.3 0.3

b = 1448.6 0.3

c = 1402.5 0.3

[1981Hig, V-C2]

labeled as Mg7Zn3 in [Mas2, 1992Aga]

340


MSIT®

Li–Mg–Zn

MgZn

< 347

oP48

?

a = 533

b = 923

c = 1716

at 52 at.% Zn [Mas2, 1992Aga]

~1 at.% of solubility range [Mas2]

[1959Ray]

determined initially as

hexagonal by [1933Tar]

Mg2Zn3

< 416

mC110

c2/m

Mg4Zn7 a = 2596

b = 524

c = 2678

= 148.6°



[1975Yar, V-C2]

labeled as Mg4Zn7 in [V-C2]

L2, MgZn2

< 586

hP12

P63/mmc

MgZn2

a = 522.3 0.1

c = 856.6 0.3

a = 521.4

c = 856.3

at 66.0-67.0 at.% Zn at 381°C [1994Goe]

[1994Goe, V-C2]

[1974Mel1]

Mg2Zn11

< 381

cP39

Pm3

Mg2Zn11

a = 855.2 0.5

at 84.1-84.6 at.% Zn at 368°C [1994Goe]

[V-C2]

* Li2MgZn cF16

Fd3m

NaTl

a = 624

a = 667 to 670

[1968Pau], metastable phase

[1950Bus, 1956Jon, 1956Wei, 1971Tes]

* Mg7(Li,Zn)9 mC192 a = 4424

b = 524

c = 1425

= 109.5°

[1974Yar], b and c axes interchanged

to get conventional monoclinic cell

composition about: Li25Mg24Zn

* L3a, Li0.77MgZn1.23 cF24

Fd3m

MgCu2

a = 744.8

an = 522.6

cn = 1290

[1974Mel1, V-C2]

constants of the equivalent

pseudohexagonal cell

* L3’, Li0.56MgZn1.44 hR* a = 1051.0

c = 1285

[1974Mel1, V-C2]

* L4’, LiMg2Zn3 hP96

LiMg2Zn3

a = 1046.0

c = 1705.0

[1974Mel1, 1974Mel2, V-C2]

* L4, Li0.25MgZn1.75 hP24

P63/mmc

MgNi2

a = 523.0

c = 1723.5

a = 522.7

c = 1709

[1958Kry]

[1974Mel1, V-C2]

* L10, Li0.23MgZn1.77 hP* a = 522.3

c = 4278

[1974Mel1, V-C2]

* L9, Li0.20MgZn1.80 hR54

R3m

Er2Co7

a = 523.0

c = 3841

a = 522.0

c = 3841

[1971Kry, V-C2]

[1974Mel1, V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

341


MSIT®

Li–Mg–Zn

a For the Laves phases polytypes, Ln, the symbol n denotes the number of superimposed layers (with hexagonal-

symmetry). The symbol Ln’ represents superstructures with doubled unit cell edge a.

* L14, Li0.11MgZn1.89 hP84

Li0.11MgZn1.89

a = 521.5

c = 5989

[1974Kry1, 1974Mel1, V-C2]

* L8, Li0.07MgZn1.93 hP* a = 521.3

c = 3422

[1974Mel1, V-C2]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%

(Mg)

βLi2Zn

3

L

(βLi)

L

Lav

es p

hase

s

Fig. 1: Li-Mg-Zn.

Partial isothermal

section at 400°C

342


MSIT®

Li–Mg–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Li Mg


Axes: at.%

(Mg)

MgZn

Mg2Zn

3

L2, MgZn

2

βLi2Zn

3

L8L

14

L9

L10

L4

L4'

L3'

L3

L

(βLi)

βLi2Zn

3+(βLi)+L

3

Fig. 2: Li-Mg-Zn.

Partial isothermal

section at 300°C

343


MSIT®

Mg–Nd–Y

Magnesium – Neodymium – Yttrium

Joachim Groebner

Literature Data

[1971Svi2] determined the solubility of Nd and Y in magnesium by microstructural examination and

electric resistivity measurements. The samples were annealed at 300, 400 and 500°C. These results of

[1971Svi2] are reported in review by [1977Ray] and [1978Dri2].

[1978Dri1] investigated the isopleth at constant 75 mass% Mg. Both papers reported a NdMg9 phase in

equilibrium with (Mg). They made no structural investigations, so the NdMg9 phase is considered to be

identical with the Nd5Mg41 phase of the binary Mg-Nd system [Mas2], also confirmed by later

investigations from [1990Del].

Binary Systems

The two binary Nd-Y and Mg-Y systems are accepted from [Mas2]. Due to latest investigations the Mg-Nd

system was accepted from [1990Del] represented in Fig. 1. For the solubility of Nd in (Mg) the values of

[1971Svi2] were preferred compared to those of [1990Del].

Solid Phases

No ternary phase has been found. All phases are listed in Table 1. The YMg1+x and NdMg phases probably

show a complete solid solution because of the identical structure and the same behavior in the La-Mg-Y

system found by [1995Gio].


One ternary invariant equilibrium is reported by [1971Svi1] and confirmed by [1978Dri1], the eutectic

L (Mg)+Y5Mg24+x+Nd5Mg41 at 536°C.

A partial reaction scheme derived from the information in [1971Svi, 1978Dri1] and the binary systems is

given in Fig. 2.

Liquidus Surface

A partial projection of the liquidus surface and the monovariant lines in the Mg-rich corner derived from

[1971Svi1] and [1978Dri1] are shown in Fig. 3.

Isothermal Sections

Figures 4 to 6 show the isothermal sections of the Mg corner at 300, 400 and 500°C after [1971Svi2].


The vertical sections at constant 75 mass% Mg, 80 mass% Mg, 12 mass% Y and 4 mass% Nd after

[1971Svi1] are reported in Figs. 7 to 10.


[1988Vos] studied hardness and electrical resistivity of annealed Mg-Nd-Y alloys. Neodymium was found

to cause a phase transformation in these alloys that leads to pronounced hardening at 200°C. The observed

precipitations could not be identified by TEM. The authors presume the formation of metastable phases, like

in binary Mg-Nd alloys, to be responsible for the hardening effect.

344


MSIT®

Mg–Nd–Y

References

[1971Svi1] Sviderskaya, Z.A., Bvidersava, Z.A., “Character of Phase Interaction in the Magnesium

Rich Part of the System Mg-Nd-Y”, Strukt. Svoistv. Legk, Splavov, 6-10 (1971) (Equi.


[1971Svi2] Sviderskaya, Z.A., Padezhnova, E.H., “Solid Solubility of Neodymium and Yttrium in

Magnesium”, Russ. Metall., 141-144 (1971), translated from Izv. Akad. Nauk SSSR, Met., 6,


[1977Ray] Raynor, G.V.,“Consitution of Ternary and Some More Complex Alloys of Magnesium”,

Int. Met. Rev., 22, 65-96 (1977) (Review, Equi. Diagram, 93)

[1978Dri1] Drits, M.E., Padezhnova, E.M., Dobatkina, T.V., “Physico-Chemical Interactions of

Elements in Magnesium Alloys of the Mg-Y-Me Systems” (in Russian), in “Magnesium

Alloys”, Akad. Nauk SSSR, Baikov Institute of Metallurgy, 74-78 (1978) (Equi. Diagram,

Experimental, 4)

[1978Dri2] Drits, M.E., Padezhnova, E.M., Dobatkina, T.V., “Effect of Additional Alloying on the

Solubility of Yttrium in Magnesium”, (in Russian), Probl. Metalloved. Tsvetn. Splavov,

Nauka, Moscow, 89-91 (1978) (Review, Equi. Diagram, 5)

[1982Gsc] Gscheidner, K.A., Jr., Calderwood, F.W., “The Nd-Y (Neodymium-Yttrium) System”, Bull.

Alloy Phase Diagrams, 3, 202-205 (1982) (Review, Equi. Diagram, 8)

[1988Vos] Vostry, P., Stulikova, I., Smola, B., Cieslar, M., Mordike, B.L., “A Study of the

Decomposition of Supersaturated Mg-Y-Nd, Mg-Y and Mg-Nd Alloys”, Z. Metallkd.,

79(5), 340-344 (1988) (Experimental, 14)

[1990Del] Delfino, S., Saccone, A., Ferro, R., “Phase Relationships in the Neodymium-Magnesium

System”, Metall. Trans. A, 21A, 2109-2114 (1990) (Experimental, Equi. Diagram, 34)

[1995Gio] Giovanini, M., Saccone, A., Marazza, R., Ferro, R., “The Isothermal Section at 500°C of the

Y-La-Mg Ternary System”, Met. Mater. Trans. A, 26A, 6-10 (1995) (Equi. Diagram,

Experimental, 28)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Mg)

<650

hP2

P63/mmc

Mg

a = 320.944

c = 521.076


[Mas2]

( YxNd1-x)

( Y)

1522 - 1478

( Nd)

1021 - 863

cI2

Im3m

W

a = 407

a = 413

0 x 1

[Mas2]

at 883°C [V-C2]

( YxNd1-x)

( Y)

< 1478

hP2

P63/mmc

Mg

a = 365.15

c = 574.74

0.4 x 1 [1982Gsc]

pure Y at 25°C [V-C2]

( ´YxNd1-x) hP4

P63/mmc

La

a = 365.62

c = 1180.56

0 x 0.3 [1982Gsc]

pure Nd at 0°C [V-C2]

In the phase diagrams of [1982Gsc,

Mas2] the and ´ phases are not

distinguished

345


MSIT®

Mg–Nd–Y

( YxNd1-x)

< 630

hR3

R3m

Sm

a = 362 to 367

c = 2620 to 2640

0.3 x 0.4 [1982Gsc, Mas2]

Nd5Mg41

< 560

tI92

I4/m

Ce5Mg41

a = 1476

c = 1039

[V-C2], considered as NdMg9 by

[1972Svi1, 1978Dri1]

NdMg3

< 744

cF16

Fm3m

BiF3

a = 739.1 [V-C2]

NdMg2

775 - 680

cF24

Fd3m

Cu2Mg

a = 866.2 [V-C2]

Y5Mg24+x

< 605

cI58

I43m

Mn

a = 1127.8

a = 1125.0



YMg2

< 780

hP12

P63/mmc

MgZn2

a = 603.7

c = 975.2

[V-C2]

(Y,Nd)Mg1+x

YMg

< 935

NdMg

750

cP2

Pm3m

CsCl a = 381.0 to 378.1

a = 386.7

probably complete solid solubility

~50 to 52 at.% Mg [V-C2]

[V-C2]

20 40 60 80

0

250

500

750

1000

Nd Mg

Mg, at.%

Te

mp

era

ture

, °C

L

800780775

545

750

650

560

NdMg NdMg3Nd5Mg41

(αNd)

(βNd)

NdMg2

Fig. 1: Mg-Nd-Y.

The binary Mg-Nd

system

346


MSIT®

Mg–Nd–Y

Fig. 2: Mg-Nd-Y. Partial reaction scheme

Mg-Nd A-B-CMg-Nd-Y Mg-Y

l (Mg)+Nd5Mg

41

545 e

l (Mg)+Y5Mg

24

566 e

L (Mg)+Nd5Mg

41+Y

5Mg

24536 E

(Mg)+Y5Mg

24+Nd

5Mg

41

L+Nd5Mg

41+Y

5Mg

24

10

10

90

Y 20.00

Nd 0.00

Mg 80.00

Y 0.00

Nd 20.00

Mg 80.00


Axes: at.%

500

566°C

600

E1, 536

(Mg)

NdMg3

Y5Mg

24

545°C

Fig. 3: Mg-Nd-Y.

Partial liquidus

surface

347


MSIT®

Mg–Nd–Y

Y 5.00

Nd 0.00

Mg 95.00

Y 0.00

Nd 5.00

Mg 95.00


Axes: at.%

(Mg)

(Mg) + NdMg12

(Mg) + NdMg12

+ Y5Mg

24+x

(Mg) + Y5Mg

24+x

Y 5.00

Nd 0.00

Mg 95.00

Y 0.00

Nd 5.00

Mg 95.00


Axes: at.%

(Mg)

(Mg) + NdMg12

(Mg) + NdMg12

+ Y5Mg

24+x

(Mg) + Y5Mg

24+x

Fig. 5: Mg-Nd-Y.


400°C

Fig. 4: Mg-Nd-Y.


300°C

348


MSIT®

Mg–Nd–Y

500

600

Y 0.00

Nd 5.32

Mg 94.68

Y 8.35

Nd 0.00

Mg 91.65Y, at.%

Te

mp

era

ture

, °C

L

L + (Mg)

(Mg) + Nd5Mg41

536°C

Y5Mg24+x

L + (Mg) + Y5Mg24+x

566°C

548°C

L + (Mg) + Nd5Mg41

(Mg) + Nd5Mg41 + Y5Mg24+x

(Mg) +

2 6

Y 5.00

Nd 0.00

Mg 95.00

Y 0.00

Nd 5.00

Mg 95.00


Axes: at.%

(Mg)

(Mg) + NdMg12

(Mg) + NdMg12

+ Y5Mg

24+x

(Mg) + Y5Mg

24+x

Fig. 7: Mg-Nd-Y.

Isopleth at

75 mass% Mg

Fig. 6: Mg-Nd-Y.


500°C

349


MSIT®

Mg–Nd–Y

500

600

700

Y 0.00

Nd 4.04

Mg 95.96

Y 6.40

Nd 0.00

Mg 93.60Y, at.%

Te

mp

era

ture

, °C

L

L + (Mg)

L + (Mg) + Nd5Mg41

L + (Mg) + Y5Mg24+x

Y5Mg24+x(Mg) + Nd5Mg41

(Mg)+Nd5Mg41+Y5Mg24+x

536°C(Mg) +

2 4 6

10

500

600

Y 0.00

Nd 0.70

Mg 99.30

Y 11.92

Nd 0.92

Mg 87.17Y, at.%

Te

mp

era

ture

, °C

L

L + (Mg)

(Mg) + Nd5Mg41

(Mg) + Y5Mg24+x

536°C

(Mg)+Nd5Mg41+

Y5Mg24+x

L + Y5Mg24+x

L + (Mg) + Y5Mg24+x

2 6

Fig. 8: Mg-Nd-Y.

Isopleth at

80 mass% Mg

Fig. 9: Mg-Nd-Y.

Isopleth at

4 mass% Nd

350


MSIT®

Mg–Nd–Y

500

600

700

Y 3.59

Nd 0.00

Mg 96.41

Y 4.71

Nd 6.29

Mg 89.00Nd, at.%

Te

mp

era

ture

, °C

L

L + (Mg)

L+NdMg3L + NdMg3 + Nd5Mg41

L + Nd5Mg41

(Mg)

L+ (Mg) + Y5Mg24+x

L + (Mg) + Nd5Mg41 536

L + Nd5Mg41 + Y5Mg24+xY5Mg24+x

(Mg) + Nd5Mg41 + Y5Mg24+x

Nd5Mg41 + Y5Mg24+x

(Mg) +

2 4 6

Fig. 10: Mg-Nd-Y.

Isopleth at

12 mass% Y

351


MSIT®

Mg–Nd–Zr

Magnesium – Neodymium – Zirconium

James Robinson, Christian Baetzner, Nathalie Lebrun, Athanasios Stamou

Literature Data

The available data on the Mg-Nd-Zr system are limited to the Mg rich part. [1978Dri] examined the (Mg)

solid solubility range at 200 and 500°C. Alloys were prepared from 99.975% pure Mg, 99.64% pure Nd and

Mg-2 mass% Zr master alloys. Some specimens were annealed at 500°C for 24 h and the remainder were

annealed for 10 h at 500°C and 200 h at 200°C, in evacuated ampoules. All the specimens were water

quenched. To determine the solid solubility of Nb and Zr in (Mg) electrical resistivity measurements and

optical metallography were carried out. Isothermal sections at 500 and 200°C are presented, showing (Zr)

and Nd5Mg41 - not NdMg12 - in equilibrium with Mg solid solution.

[1980Zak] presents a partial liquidus surface giving no experimental details. He states the presence of a

ZrMg2 phase in equilibrium with (Mg) and the Mg-Nd phases, which is inconsistent with the results of

[1978Dri].

Binary Systems

The Mg-Zr phase diagram is accepted from [Mas2], the Mg-Nd system, however from [1990Del] since

[Mas2] gives only a tentative diagram. The ZrMg2 phase found by [1980Zak] is not accepted in this

evaluation since [Mas2] states for the binary system that this phase is stabilized by oxygen and other

impurities. Besides that it has not been found in the ternary experimental work of [1978Dri] in the Mg

corner, where (Mg) is in equilibrium with (Zr). Compared to the version of [Mas2] in the diagram of

[1990Del] for the Mg-Nd system the phase NdMg12 does not appear. It is stated to be metastable and only

exists after quenching from the liquid state. This is in agreement with the results for the ternary system from

[1978Dri, 1980Zak], and therefore accepted here.

Solid Phases

Based on the available data no ternary phases exist in this system. Crystal structure data on the phases within

the known part of the system are presented in Table 1.

[1978Dri] presented tie-line triangles which seem to indicate solubility of Zr in the Mg-Nd phases, but no

quantitative experimental evidence is given.


[1980Zak] states the ZrMg2-NdMg3 section to be a pseudobinary eutectic, this is not accepted based on the

binary Mg-Zr system. Since there are no intermediate phases known in the Mg-Zr system, a pseudobinary

section is introduced between (Zr) and NdMg3.


The invariant equilibria deduced from the liquidus projection from [1980Zak] are presented in the reaction

scheme given in Fig. 1. [1980Zak] gave an estimated temperature for U1 of 600°C and a binary peritectic

p2 temperature of 640°C. This temperature has changed to 560°C in the Mg-Nd version of [1990Del] so the

temperature of the ternary reaction U1 is now unknown, as well as U2, which has been given ~580°C by

[1980Zak], which again is inconsistent with the reaction sequence and the binary peritectic p2. Therefore

no temperatures can be given for U1 and U2.

Liquidus Surface

The presented liquidus surface is shown in Fig. 2 and is based on the work of [1980Zak]. The binary

invariant reactions have been shifted in accordance with [Mas2, 1990Del]. [1980Zak] has found the ZrMg2

352


MSIT®

Mg–Nd–Zr

phase which is stabilized by impurities [1988Nay], thus putting some doubt onto the remainder of the

diagram. The data from [Mas2] for Mg-Nd is tentative. For example the position of e2 ranges from 1.7 at.%

to 5.6 at.% Nd, whereas [1990Del] gives 7.5 at.% Nd.

The positions of the ternary invariant reactions have also been shifted, compared to [1980Zak], in

accordance with the binary invariant reactions. This, of course, gives the complete liquidus surface and the

reaction scheme a tentative character.

Further experimental work is necessary to confirm the entrance of U1 and U2 as well as to determine the

primary crystallization areas in this region of the diagram.

Isothermal Sections

Partial isothermal sections at 500°C and 200°C are presented in Fig. 3 and 4 based on the work of [1978Dri].

The phase notation NdMg9 of [1978Dri] has been replaced by Nd5Mg41.

References

[1978Dri] Drits, M.E., Padezhnova, E.M., Guzei, L.S., “The Magnesium-Neodymium-Zirconium

Phase Diagram (Magnesium Rich Range)”, Russ. Metall., 1, 195-198, (1978), translated

from Izv. Akad. Nauk SSSR, Met., 1, 218-220, (1978) (Experimental, Equi. Diagram, #, 13)

[1980Zak] Zakharov, A.A., “Promyshlennye Splavy Tsvetnykh Metallov” (in Russian), 108-109,

(1980) (Review, Equi. Diagram, #, 259)

[1988Nay] Nayeb-Hashemi, A.A., Clark, J.B., “Mg-Zr (Magnesium-Zirconium)”, in “Phase Diagrams

of Binary Magnesium Alloys”, ASM International, Metals Park, Ohio, 365-370 (1988)


[1990Del] Delfino, S., Saconne, A., Ferro, R., “Phase Relationships in the Neodymium-Magnesium

Alloy System”, Metall. Trans. A., 21A, 2109-2114, (1990) (Experimental, Equi.

Diagram, #, 34)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Nd)(h)

1021 - 863

cI2

Im3m

W

a = 413 [Mas2]

( Nd)(r)

< 863

hP4

P63/mmc

La

a = 365.82

c = 1179.66

at 25°C [Mas2]

( Zr)(h)

1855 - 863

cI2

Im3m

W

a = 360.90 [Mas2]

( Zr)(r)

< 863

hP2

P63/mmc

Mg

a = 323.16

c = 514.75

at 25°C [Mas2]

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.944

c = 521.076

at 25°C [V-C2]

[Mas2]

353


MSIT®

Mg–Nd–Zr

NdMg12

< 591

tI26

I4/mmm

Mn12Th

a = 1031

c = 593

[V-C2]

metastable phase [1990Del, 1978Dri,

1980Zak]

Nd5Mg41

< 560

tI92

I4/m

Ce5Mg41

a = 1476

c = 1039

a = 1474.1

c = 1039.6

[V-C2]

[1990Del]

NdMg3

< 780

cF16

Fm3m

BiF3

a = 739.1 [1990Del]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

Fig. 1: Mg-Nd-Zr. Reaction scheme

Mg-Zr A-B-CMg-Nd-Zr Mg-Nd

L (αZr)+NdMg3

>600 e1

l + (αZr) (Mg)

653.6 p1 ?

l+NdMg3

Nd5Mg

41

560 p2

L+NdMg3

(αZr)+Nd5Mg

41<560 U

1

l (Mg)+Nd5Mg

41

545 e2

L+(αZr)+Nd5Mg

41

L+(αZr) (Mg)+Nd5Mg

41>545 U

2

(αZr)+(Mg)+Nd5Mg

41

(αZr)+Nd5Mg

41+NdMg

3

354


MSIT®

Mg–Nd–Zr

20

40

60

80

20 40 60 80

20

40

60

80

Nd Zr


Axes: at.%

(αZr)

(Mg)

NdMg3

e1

U1

U2

Nd5Mg

41

NdMg3

p1, 653.6°C

e2, 545°C

p2, 560°C

Fig. 2: Mg-Nd-Zr.

Schematic partial

liquidus surface

Nd 1.00

Zr 0.00

Mg 99.00

Nd 0.00

Zr 1.00

Mg 99.00


Axes: at.%

(Mg)

(Mg) + (αZr)

(Mg) + (αZr) + Nd5Mg

41

(Mg) + Nd5Mg

41

Fig. 3: Mg-Nd-Zr.

Partial isothermal

section at 500°C

355


MSIT®

Mg–Nd–Zr

Nd 0.20

Zr 0.00

Mg 99.80

Nd 0.00

Zr 0.20

Mg 99.80


Axes: at.%

(Mg)

(Mg) + (αZr)(Mg) + Nd

5Mg

41

(Mg) + (αZr) + Nd5Mg

41

Fig. 4: Mg-Nd-Zr.

Partial isothermal

section at 200°C

356


MSIT®

Mg–Sn–Zn

Magnesium – Tin – Zinc

Lazar Rokhlin

Literature Data

[1933Ota] studied the Mg-Sn-Zn phase diagram using thermal analysis and microscopy method. The phase

relations in solid state and the phase reactions during solidification were investigated in the system in the

full concentration range. The investigation showed existence of two pseudobinary sections and a number of

the invariant four-phase reactions in the system. However, [1933Ota] used the old view of the Mg-Zn binary

phase diagram without some significant details recognized later. So, the Mg-Zn phase diagram used by

[1933Ota] showed four intermediate solid phases, (at approximately 30 at.% Zn), MgZn, MgZn2 and

MgZn5 unlike the modern view of the Mg-Zn phase diagram with five intermediate solid phases, Mg7Zn3,

MgZn, Mg2Zn3, MgZn2 and Mg2Zn11. Moreover, there are differences in invariant reactions between the

modern Mg-Zn phase diagram and that had used by [1933Ota].

[1959Gla] determined the solubility of tin in the phase MgZn2 at one temperature 400°C using the X-ray

method.

[1986Min] studied the structure of the Mg-Sn-Zn cast alloys in as cast condition by the Moessbauer

spectroscopy and X-ray diffraction methods. The alloy compositions corresponded to the formulae

(Mg70Zn30)100-xSnx, where x = 2.5, 5, 7.5, 10, 12.5 or 15. In all studied alloys [1986Min] found tin to be

only associated with Mg into Mg2Sn. This fact suggested the very small solubility of tin the Mg-Zn phases

being in equilibrium with Mg2Sn. In the structure of some alloys [1986Min] observed Mg2Sn together with

phases Mg7Zn3 and Mg2Zn3 suggesting a possibility of equilibria between Mg2Sn and these binary phases

of the Mg-Zn system.

[1987Sir] presented actually results of the same investigation of the cast Mg-Sn-Zn alloys as [1986Min].

The main conclusions of [1987Sir] concerning the Mg-Sn-Zn phase diagram coincided with those of

[1986Min].

[1989Min] conducted some additional experiments to the [1986Min, 1987Sir] studies on the Mg-Sn-Zn

alloys and confirmed in general the results of [1986Min, 1987Sir] on the Mg-Sn-Zn phase diagram.

[1994Goe] studied the Mg-Sn-Zn phase diagram in the Mg2Sn-MgZn2-Zn-Sn area by the thermal analysis

and microscopy method. This work was quite detailed and based on the modern versions of the binary phase

diagrams. Existence of the same pseudobinary sections and the same four-phase invariant reactions

established by [1933Ota] earlier have been confirmed by [1994Goe], but their temperatures and the

compositions of the liquid phase have been refined. Besides, [1994Goe] took into account the accepted

versions of the adjoining binary phase diagrams. [1994Goe] confirmed the sequence of the invariant

reactions in the Mg2Sn-MgZn2-Zn-Sn area established by [1933Ota]. In [1994Goe] a number of the

polythermal vertical sections of the phase diagram and the projection of the liquid surface for the

Mg2Sn-MgZn2-Zn-Sn area were constructed. Two pseudobinary sections of the phase diagram were

constructed as well.

Binary Systems

The phase diagrams of the binary systems Mg-Sn, Mg-Zn and Sn-Zn are accepted from [Mas2].

Solid Phases

No ternary phase has been found in the system. The solid unary and binary phases are listed in Table 1. Solid

Mg dissolves up to 14.48 mass% Sn (3.35 at.% Sn) and up to 6.2 mass% Zn (2.4 at.% Zn) at the eutectic

temperatures 561.2 and 340°C, respectively [Mas2]. Extension of the Sn- and Zn-base solid solution is quite

insignificant [Mas2]. The most well established values of the solubility of other components in these solid

solutions are 0.15 mass% Mg (0.4 at.% Mg) in solid Zn [Mas2] and about 0.3 mass% Zn (0.6 at.%Zn) in

solid Sn [Mas2]. Solubility of Zn in Mg2Sn along the Mg2Sn-MgZn2 section reaches about 0.2 mass% and

357


MSIT®

Mg–Sn–Zn

about 0.1 mass% along the Mg2Sn-Zn section [1994Goe]. [1933Ota] showed higher solubility of Zn in

Mg2Sn (up to about 5 mass%), but these data are not convincing. There is no solubility of Sn in Mg7Zn3

and Mg2Zn3 [1986Min, 1987Sir]. The solubility of Sn in MgZn2 is 3.6 mass% (1.6 at.%) with constant Mg

content at 400°C according to [1959Gla] and 0.6 mass% Sn along the section Mg2Sn-MgZn2 at 340°C

according to [1994Goe]. The data of [1994Goe] are preferable as they have been obtained basing on the

more detailed investigation. The maximum solubility of Sn in MgZn2 along the section Mg2Sn-MgZn2

reaches 0.7 mass% at the eutectic temperature 567°C [1994Goe]. Solubility of Sn in other Mg-Zn

intermediate phases (MgZn, Mg2Zn11) was not studied.


The sections Zn-Mg2Sn and MgZn2-Mg2Sn are established by [1933Ota] and confirmed by [1994Goe] as

pseudobinary systems. The pseudobinary systems are displayed in Figs. 1 and 2. They are drawn according

to [1994Goe] assuming the data of this work to be more reliable as compared with those of the earlier

investigation [1933Ota]. The solubility of Sn in MgZn2 in the system MgZn2-Mg2Sn (Fig. 2) determined

by [1959Gla] is also rejected as less reliable than that presented by [1994Goe].


The ternary invariant equilibria in the Mg-Sn-Zn system are presented in Table 2. The invariant equilibria

out of the Mg2Sn-MgZn2-Zn-Sn area (e7, e8, U3, E2,) are accepted mainly according to [1933Ota], but with

some corrections to meet the modern versions of the Mg-Zn and Mg-Sn phase diagrams. The invariant

equilibria within the area Mg2Sn-MgZn2-Zn-Sn are accepted from [1994Goe] recognizing the data of

[1994Goe] to be more reliable and compatible with the boundary systems than those of [1933Ota] for the

same reactions. Two invariant four-phase equilibria have to be supposed to meet the sequence of the

intermediate phase formation during solidification in the Mg-Zn system. These invariant equilibria are

L + MgZn2 Mg2Zn3 + Mg2Sn (U1) and L + MgZn Mg7Zn3 + Mg2Sn (U4). The types and temperatures

of these reactions are shown as tentative taking into account the Mg-Zn phase diagram and the location of

the monovariant reaction lines shown by [1933Ota]. Compatibility with the accepted Mg-Zn phase diagram

has required also to replace the four-phase equilibrium L + MgZn2 MgZn + Mg2Sn [1933Ota] with the

equilibrium L + Mg2Zn3 MgZn + Mg2Sn (U3) assuming the same composition of the liquid phase.

Moreover, it is reasonable to suppose in the system a four-phase equilibrium related to the decomposition

of Mg7Zn3 by the eutectoid reaction e6 in the binary Mg-Zn system. Only solid phases have to take part in

this equilibrium. As far as the solubility of Sn in Mg7Zn3 is actually absent it is reasonable to assume this

equilibrium to be of the degenerate type with its temperature coinciding with the temperature of the

Mg7Zn3 (Mg) + MgZn equilibrium in the Mg-Zn binary system.

The reaction scheme is presented in Fig. 3, and projection of the invariant equilibrium planes with

connecting lines of double saturation are presented in Fig. 4. They were constructed basing on the data

[1933Ota] out of the Mg2Sn-MgZn2-Zn-Sn area and on the data [1994Goe] within the

Mg2Sn-MgZn2-Zn-Sn area. Some invariant equilibrium temperatures and dispositions of the points and

lines were corrected to make compatible the accepted boundary systems and the results of the investigations

[1933Ota] and [1994Goe] for the fields out of and within the Mg2Sn-MgZn2-Zn-Sn area, respectively.

Thus, aiming to meet the accepted Mg-Zn phase diagram the temperatures of the equilibria U3 and E2 were

assumed to be 346 and 339°C, respectively, as compared with 351 and 340°C presented by [1933Ota]. By

the same way the temperatures of the supposed equilibria U1 and U4 were assumed, as well. The

temperatures of the other equilibria were taken according to [1994Goe] without changes.

Compositions of the solid phases taking part in the invariant equilibria are also corrected taking into account

the accepted binary phase diagrams.

Liquidus Surface

Figure 5 displays the isotherms of the liquidus surface and melting grooves separating the fields of primary

crystallization. The primary crystallization fields of (Mg), Mg2Sn, MgZn2, (Zn), ( Sn) are marked on the

358


MSIT®

Mg–Sn–Zn

plot. The remaining primary crystallization fields are e3E2U4p4 for Mg7Zn7, p4U4U3p3 for MgZn,

p3U3U1p1 for Mg2Zn3 and p2U2E1e2 for Mg2Zn11. Positions of the pseudobinary sections Zn-Mg2Sn and

MgZn2-Mg2Sn are shown by the dashed lines in Fig. 5.

Isothermal Sections

Figure 6 displays the isothermal section at about 250°C constructed using the data [1933Ota, 1994Goe] with

corrections and additions to meet the boundary systems and the results of the investigations [1986Min,

1987Sir]. The magnesium solid solution field was delineated taking into account the solubilities of Sn and

Zn in solid Mg in the binary systems at 250°C. Because both solubilities are small it is assumed that the Sn

and Zn atoms does not interact in Mg solid solution and, therefore, the solubility of each of them in solid

Mg in ternary system is assumed to be similar that in the respective binary system. Following [1986Min,

1987Sir] and the accepted Mg-Zn phase diagram the existence of the areas with the phase Mg2Zn3

(Mg2Sn+MgZn+Mg2Zn3, Mg2Sn+Mg2Zn3, Mg2Sn+Mg2Zn3+MgZn2) is assumed.


Three polythermal vertical sections of the phase diagram are presented in Figs. 7-9. The sections for 40 and

85 mass% Sn (Figs. 7 and 8) were constructed following [1994Goe]. The section for 10 mass% Sn (Fig. 9)

was constructed following [1933Ota] within the Mg2Sn-MgZn2-Zn-Sn concentration area and following

[1994Goe] out of this concentration area. Some corrections were made also to meet the accepted binary

phase diagrams.

References

[1933Ota] Otani, B., “Constitution of the Phase Equilibrium Diagram of the Magnesium-Zinc-Tin

System” (in Japanese), Tetsu to Hagane, 19, 566-574 (1933) (Experimental, Equi. Diagram,

#, 6)

[1933Tar] Tarschisen, L., “X-ray Investigation of the MgZn and MgZn5 Compounds”,

Z. Kristallogr. A, A186, 423-438 (1933) (Crys. Structure, 6)

[1959Gla] Gladyshevsky, E.I., Cherkashin, E.E., “Solid Solutions Based on Metallic Compounds” (in

Russian), Zh. Neorg. Khim., 1(6), 1394-1401 (1959) (Experimental, Equi. Diagram, 4)


Magnesium and its Alloys”, Pergamon Press, London, New York, Paris, Los Angeles,


[1975Yar] Yarmolyuk, Ya.P., Kripiakevich, P.I., Melnik, E.V., “The Crystal Structure of the Mg4Zn7

Compound”, Sov. Crystallogr., 20(3), 538-542 (1975), translated from Kristallografiya,

20(3), 538-542 (1975), (Experimental, Crys. Structure, 16)



[1986Min] Mingolo, M., Arcondo, B., Nassif, E., Sirkin, H., “Changes in the Glass Forming Ability of

MgZnSn Alloys Due to the Presence of an Intermetallic Compound”, Z. Naturforsch. A,

A41(12), 1357-1360 (1986) (Experimental, Equi. Diagram, 14)

[1987Sir] Sirkin, H., Mingolo, N., Nassif, E., Arcondo, B., “Increase of the Glass-Forming

Composition Range of Mg-Based Binary Alloys by Addition of Tin”, J. Non-Cryst. Solids,

93(2-3), 323-330 (1987) (Experimental, Equi. Diagram, 16)

[1989Min] Mingolo, N., Nassif, E., Arcondo, B., Sirkin, H., “Two Competitive Effects in the

Glass-Forming Ability of Mg-Based Alloys”, J. Non-Cryst. Solids, 113(2-3), 161-166





359


MSIT®

Mg–Sn–Zn





Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.944

c = 521.076

a = 320.99 to

319.57

c = 521.08 to

518.82

a = 320.68

c = 521.49

pure Mg at 25°C [V-C2, Mas2]


at 0 to 2.81 at.% Zn [V-C2]

dissolves 3.35 at.% Sn at 561°C [Mas2]

at 2.58 at.% Sn at 25°C [V-C2]

( Sn)

231.9681-13

tI4

I41/amd

Sn

a = 583.18

c = 318.18

pure Sn at 25°C [Mas2]

dissolves up to 0.6 at.% Mg at 204°C

and 0.6 at.% Zn at 199°C [Mas2]

Sn

< 13

cF8

Fd4m

C(diamond)

a = 648.92 [Mas2]

(Zn)

< 419.5

hP2

P63/mmc

Mg

a = 266.47

c = 494.69


dissolves up to 0.4 at.% Mg at 364°C

and 0.039 at.% Sn at 199°C [Mas2]

Mg2Sn

< 770.5

cF12

Fm3m

CaF2

a = 676.5 [V-C2, Mas2]

Mg7Zn3

342-325

oI158

Immm

Mg51Zn20

a = 1408.3 0.3

b = 1448.6 0.3

c = 1402.5 0.3

[1981Hig, V-C2]

refined as Mg51Zn20

MgZn

< 347

oP48

?

a = 533

b = 923

c = 1716

52 at.% Zn [Mas2, 1992Aga]


[1959Ray]



Mg2Zn3

< 416

mC110

C2/m

Mg4Zn7 a = 2596

b = 524

c = 2678

= 148.6°

60 at.% Zn [Mas2, 1992Aga]


[1975Yar, V-C2]

labelled as Mg4Zn7 in [V-C2]

360


MSIT®

Mg–Sn–Zn


Parentheses < > indicate estimated values

MgZn2

< 586

hP12

P63/mmc

MgZn2

a = 522.3 0.1

c = 856.6 0.3

66.0-67.0 at.% Zn at 381°C [1994Goe]

[V-C2]

Mg2Zn11

< 381

cP39

Pm3

Mg2Zn11

a = 855.2 0.5

84.1-84.6 at.% Zn at 368°C [1994Goe]

[V-C2]

Reaction T [°C] Type Phase Composition [at.%]

Mg Sn Zn

L MgZn2 + Mg2Sn 567 e7 (max) L

MgZn2

Mg2Sn

38.59

33.46

~66.55

4.83

0.3

~33.28

56.58

66.24

~0.17

L + MgZn2 Mg2Zn3 +

Mg2Sn

<414> U1 L

MgZn2

Mg2Zn3

Mg2Sn

<64.10>

32.37

40.06

~66.71

<1.68>

0.71

0

~33.29

<34.22>

66.92

59.94

0

L + MgZn2 Mg2Zn11 +

Mg2Sn

368 U2 L

MgZn2

Mg2Zn11

Mg2Sn

9.69

~33.45

~15.32

~66.69

3.9

~0.29

~0.05

~33.22

86.41

~66.26

~84.63

~0.09

L Mg2Sn + (Zn) 355 e8 (max) L

Mg2Sn

(Zn)

8.94

~66.69

~0.4

4.15

~33.22

0

86.91

~0.09

~99.60

L Mg2Zn11 + Mg2Sn + (Zn) 353 E1 L

Mg2Zn11

Mg2Sn

(Zn)

9.17

~15.32

~66.69

0.4

3.92

0.05

33.22

0

86.91

84.63

0.09

99.60

L + Mg2Zn3 MgZn +

Mg2Sn

346 U3 L

Mg2Zn3

MgZn

Mg2Sn

69.11

40.66

47.94

66.71

0.63

0

0

33.29

30.26

59.34

52.06

0

L + MgZn Mg7Zn3 + Mg2Sn <341> U4 L

MgZn

Mg7Zn3

Mg2Sn

70.97

48.59

69.96

66.71

0.31

0

0

33.29

28.72

51.41

30.04

0

L (Mg) + Mg7Zn3 + Mg2Sn 339 E2 L

(Mg)

Mg7Zn3

Mg2Sn

71.54

97.01

69.96

66.71

0.15

0.54

0

33.29

28.31

2.45

30.04

0

L Mg2Sn + ( Sn) + (Zn) 183 E3 L

Mg2Sn

( Sn)

(Zn)

7.39

66.69

99.89

0.4

81.14

~33.22

0

0

11.47

~0.09

0.11

99.6

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

361


MSIT®

Mg–Sn–Zn

20 40 60 80

300

400

500

600

700

800

Mg 66.67

Zn 0.00

Sn 33.33

Zn

Zn, at.%

Te

mp

era

ture

, °C

~0.1

L

Mg2Sn+(Zn)

355

419.58°C

770.5°C

L+(Zn)

L+Mg2Sn

88.9

(Zn)

Mg2Sn

Fig. 1: Mg-Sn-Zn.

The pseudobinary

system Zn - Mg2Sn

20 40 60

300

400

500

600

700

800

Mg 66.67

Zn 0.00

Sn 33.33

Mg 33.33

Zn 66.67

Sn 0.00Zn, at.%

Te

mp

era

ture

, °C

770.5°C

590°C

L+Mg2Sn

~0.2 0.771

567

L+MgZn2

Mg2Sn+MgZn2

L

Mg2Sn MgZn2

Fig. 2: Mg-Sn-Zn

The pseudobinary

system

MgZn2 - Mg2Sn

362


MSIT®

Mg–Sn–Zn

Fig

. 3:

Mg-

Sn-Z

n. R

eact

ion s

chem

e

Sn

-Zn

Mg

-Sn

Mg-Z

nM

g-S

n-Z

n

l (

βSn)

+ (

Zn)

198.5

e 6

l (

Mg)

+ M

g2S

n

561.2

e 1

l M

g2S

n +

(βS

n)

203.5

e 5

l +

MgZ

n2

Mg

2Z

n3

41

6p

1

l +

MgZ

n2

Mg

2Z

n1

1

38

1p

2

l M

g2Z

n1

1 +

(Z

n)

36

4e 4

l +

Mg

2Z

n3

MgZ

n

34

7p

3

l (

Mg)

+ M

g7Z

n3

34

0e 3

l +

MgZ

n

Mg

7Z

n3

34

2p

4

Mg

7Z

n3

(Mg

)+M

gZ

n

32

5e 4

L M

gZ

n2 +

Mg

2S

n

57

0e 7

L+

MgZ

n2

Mg

2Z

n3+

Mg

2S

n<

414

>U

1

L+

MgZ

n2

Mg

2Z

n1

1+

Mg

2S

n3

68

U2

L M

g2S

n +

(Z

n)

35

5e 8

L M

g2Z

n1

1+

Mg

2S

n+

(Zn)

35

3E

1

L+

Mg

2Z

n3

MgZ

n+

Mg

2S

n3

46

U3

L+

MgZ

nM

g7Z

n3+

Mg

2S

n<

341

>U

4

L (

Mg)+

Mg

7Z

n3+

Mg

2S

n3

39

E2

Mg

7Z

n3

(Mg

)+M

gZ

n,M

g2S

n3

25

D

L M

g2S

n+

(βS

n)+

(Zn)

18

3E

3L+

Mg

2Z

n1

1+

Mg

2S

n

MgZ

n2+

Mg

2Z

n3+

Mg

2S

n

MgZ

n2+

Mg

2Z

n1

1+

Mg

2S

n

Mg

2Z

n1

1+

Mg

2S

n+

(Zn)

Mg

2Z

n3+

MgZ

n+

Mg

2S

n

Mg

Zn

+M

g7Z

n3+

Mg

2S

n

(Mg

)+M

gZ

n+

Mg

2S

n

Mg

2S

n+

(βS

n)+

(Zn)

L+

Mg

2Z

n3+

Mg

2S

n

L+

MgZ

n+

Mg

2S

n

L+

Mg

7Z

n3+

Mg

2S

n

(Mg

)+M

g7Z

n3+

Mg

2S

n

363


MSIT®

Mg–Sn–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Zn


Axes: at.%

E3

e6

e5

Mg2Sn

(Mg)

e1

e2

p2U

2

e8 E

1

p1

U4

MgZn2

e7

U1

U3

e3p

4 p3

250

750

300

700

250

350

350

300

400

550

550

650

600

650600

550

(βSn)

(Zn)

400

500

500450

Mg2Sn

600

500450

E2

20

40

60

80

20 40 60 80

20

40

60

80

Mg Zn


Axes: at.%

E3

e6

e5

Mg2Sn

(Mg)

e1

e2

p2Mg

2Zn

11

U2

e8 E

1

p1

U4

MgZn2MgZn Mg

2Zn

3

e7

U1U

3

Mg7Zn3

e3

p4

p3E

2

Fig. 5: Mg-Sn-Zn.

Liquidus surface

Fig. 4: Mg-Sn-Zn.

Projection of

four-phase equilibria

planes and connected

lines of double

saturation

364


MSIT®

Mg–Sn–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Mg Zn


Axes: at.%

Mg2Sn+(βSn)+(Zn)

(Mg)

Mg2Sn

Mg2Zn

11MgZn

2MgZn Mg

2Zn

3

(Mg)+MgZn+Mg2Sn

(Zn)

40 50 60 70

100

200

300

400

500

600

700

Mg 51.77

Zn 30.10

Sn 18.12

Mg 0.00

Zn 73.14

Sn 26.86Zn, at.%

Te

mp

era

ture

, °C

L

Mg2Sn+(βSn)+(Zn)

L+Mg2Sn+(Zn)

183

L+Mg2Sn+MgZn2

L+(Zn)

L+(βSn)+(Zn)

Mg2Sn+MgZn2+Mg2Zn11

353

368

L+Mg2Sn

L+Mg2Sn+Mg2Zn11

Mg2Sn+Mg2Zn11+(Zn)

Fig. 6: Mg-Sn-Zn.


250°C

Fig. 7: Mg-Sn-Zn.


40 mass% Sn

365


MSIT®

Mg–Sn–Zn

10 20

250

500

Mg 46.29

Zn 0.00

Sn 53.71

Mg 0.00

Zn 24.26

Sn 75.74Zn, at.%

Te

mp

era

ture

, °C

L

L+(βSn)+(Zn)

L+(Zn)

Mg2Sn+(βSn)+(Zn)

L+Mg2Sn+(Zn)

L+Mg2Sn

183

L+(βSn)+(Zn)

20 40 60 80

100

200

300

400

500

600

700

Mg 97.78

Zn 0.00

Sn 2.22

Mg 0.00

Zn 94.23

Sn 5.77Zn, at.%

Te

mp

era

ture

, °C

L

Mg7Zn3+MgZn+Mg2Sn

(Mg)

325

339

L+Mg2Sn

340

414L+(Mg)+

346

L+MgZn2+

353

L+Mg2Zn3+

368

183

(βSn)+(Zn)

L+(βSn)

Mg2Sn+

L+

(Zn)

MgZn+Mg2Zn3+Mg2Sn

L+Mg2Sn

Mg2Zn11+Mg2Sn+(Zn)

L+Mg2Sn

L+Mg2Zn11+

L+MgZn2

Mg2Sn

L+MgZn2+

MgZn2+

Mg2Zn3+MgZn2+Mg2Sn

L+(Mg)

(Mg)+Mg2Sn

(Mg)+MgZn+Mg2Sn

(Mg)+Mg7Zn3+Mg2Sn

Mg2Sn

Mg2Sn Mg2Sn

Mg2Sn

+(Zn)

Mg2Zn11+

Mg2Sn

+(Zn)

Fig. 8: Mg-Sn-Zn.


85 mass% Sn

Fig. 9: Mg-Sn-Zn.


10 mass% Sn

366


MSIT®

Mg–Y–Zn

Magnesium – Yttrium – Zinc

Nathalie Lebrun, Athanasios Stamou,Christian Baetzner, James Robinson, Alexander Pisch

Literature Data

[1974Dri] mentioned the existence of three ternary phases designated W(Y3Mg2Zn3), Z(YMg3Zn6),

X(YMg12Zn). This is in agreement with other experimental works of [1977Dri, 1979Dob, 1979Pad,

1982Pad]. Formation and crystal structure of the ternary phases have been studied also recently. [1993Luo]

was the first who reported the existence of a stable icosahedral ternary phase in the Mg-Y-Zn system. Its

existence was later confirmed [1994Nii, 1994Tsa, 1995Tsa, 1997Tsa, 1998Fis, 1998Lan, 1999Abe1,

1999Abe2, 2000Ste, 2000Tsa, 2001Yi] using X-ray diffraction, high resolution TEM and calorimetric

techniques. Its X-ray diffraction pattern measured by [1994Nii, 1995Tsa] is identical to the one of the

Z phase detected by [1982Pad] although a slight difference in composition is observed. Consequently, this

icosahedral phase shows a composition close to Z(YMg3Zn6) and is called 2´ in this assessment. Two

superstructures of a hexagonal 2 were found with an identical lattice parameter c but different parameter a

[1999Abe1]. An additional decagonal quasi-crystalline phase 2´´ was also found [1998Sat, 2001Sat].

Another ternary phase has been recently reported [1999Abe1, 1999Abe2] and called 3 in this assessment.

[2000Luo] has established crystal structure of a ternary phase YMg12Zn [1977Dri] by high-resolution

electron microscopy. This phase is designated 1 in this assessment.

The primary solidification of the icosahedral 2´ phase was extensively studied by [1997Lan1, 1997Lan2].

[2000Ste, 2001Yi] also investigated its solidification behavior from Mg-Y-Zn melts with low Y content (0

to 3 at.% Y).

Reviewing previous results [1977Dri, 1977Ray] presents liquidus isotherms in the Mg-rich part. Additional

experimental works on the liquidus were done later [2001Yi] and a liquidus phase boundary (Mg)/ 2´ was

established. [1979Pad] also investigated a partial liquidus projection by microscopical examination and

differential thermal analysis of slow-cooled alloys.

[1968Zas] investigated the fusion diagram and the solid solubility of yttrium and zinc in the solid

magnesium rich region. Materials were annealed for 240 to 2280 hours from 200 to 500°C followed by

quenching in water. 36 alloys were investigated with compositions distributed along five sections of the

concentration triangle with constant Y:Zn ratios of 9:1, 7:3, 1:1, 3:7 and 1:9. Experimental results were

based on X-ray structural and thermal analyses.

[1979Dob, 1982Pad] constructed isothermal sections using differential thermal analysis, X-ray

crystallography, microstructure and electron microprobe methods and microscopical examination. Samples

were annealed for 50 hours at 500°C and 500 hours at 300°C [1979Dob]. [1982Pad] established a partial

isothermal section at 300°C ranging up to 70 at.% Zn and 50 at.% Y. [1979Dob] determined isothermal

sections at 300 and 500°C up to 20 mass% Y and 30 mass% Zn. The isothermal section at 300°C from

[1979Dob] turned out to be in good agreement with the section from [1982Pad]. Recently, [2000Tsa]

studied the equilibrium phase diagram in the region of the quasicrystals formation and established the

equilibria between 2 and 2=, and other stables phases. Isothermal sections at 427, 500 and 600°C in the

composition range (30-70)Zn-(20-60)Mg-(0-20)Y (at.%) were then established from fifteen annealed alloys

using X-ray diffraction, SEM and TEM techniques.

Several polythermal sections have been reported by [1979Pad] (constant 18 mass% Y and constant 29

mass% Zn) and [1977Dri] going from Mg - 17 %Y to Mg - 30 % Zn. A polythermal section Mg - 1 is given

by [1982Pad] and proved to be pseudobinary. An additional pseudobinary section Mg - 4 may exist

according to the results of [1979Pad].

Binary Systems

Assessment of the Mg-Y system by [2003Fab] and Mg-Zn by [2001Shc] are accepted. They are based on

[1998Luk, 1997Fla] for Mg-Y and [1992Aga, 1992Luk, 1994Goe] for Mg-Zn. The binary system Y-Zn is

accepted from [Mas2].

367


MSIT®

Mg–Y–Zn

Solid Phases

Crystallographic data for all the solid phases are presented in Table 1.

[1982Pad] states the existence of a continuous solid solution between the binary phases YMg and YZn, and

gives an intermediate lattice parameters for the solid solution, see Table 1.

Existence of three ternary compounds in this system called here 1, 2´ and 4 is assumed according to

[1982Pad]. Isothermal sections from [1979Dob] at 300 and 500°C and from [1982Pad] at 300°C seem to

indicate that all these ternary phases have a certain homogeneity ranges although the exact ranges are

unknown. [1982Pad] gives formulas for the ternary phases. However the formula assigned by [1982Pad] to

1 was altered from YMg12Zn to Y9Mg85Zn6. This was to ensure that the pseudobinary section from (Mg)

to 1 crosses the liquid monovariant curve given by [1979Pad] between e2 and E1, see Fig. 2. The

composition of 1 was chosen since [1977Dri] and [1982Pad] state that the alloy 62.7Mg-25.4Y-11.9Zn

(mass%) was practically single-phase. Further experimental work is necessary to determine the accurate

homogeneity ranges of 1 and 4 in order to check the tie-lines with (Mg) solid solution.

A large number of ideal compositions were found for 2´ in the composition range

(0.5-10)Y-(43-65)Zn-(25-48)Mg (at.%) [1994Nii, 1994Tsa, 1995Kon, 1995Tsa, 1997Lan2, 1997Tsa,

1998Fis, 1998Lan, 1999Abe2, 2000Ste]. Nevertheless, the larger quasi-crystals are obtained for a

composition close to YMg3Zn6 [1998Lan]. Consequently, the crystallization of 2´ strongly depends on the

initial composition of the alloys. [1995Kon] prepared the icosahedral phase over a wide composition range.

The primary crystallization of this phase has been precisely investigated by [1997Lan1] using DSC

technique for more than 40 alloys. The crystallization region was located between 36 and 75 at.% Mg and

up to 4 at.% Y. However discrepancies is observed since [1997Tsa] and [2001Yi] found a primarily

crystallization of 2´ in a Y8Mg42Zn50 and Y0.03Mg73.97Zn25.99 alloys outside the primarily crystallization

region found by [1997Lan1, 1997Lan2]. These discrepancies may be explained by the presence of

numerous metastable and stable quasi-crystalline phases. Consequently, no conclusion can be made

concerning the primary crystallization of 2´.

The structure of 1 has been determined by [2000Luo]. It can be indexed either by using a hexagonal or a

trigonal structure. At small compositional change two superstructures, 2´ and 2´´, of the hexagonal 2 are

characterized by an icosahedral and a decagonal structures, respectively [1993Luo, 1998Sat, 1999Abe2,

2000Tsa]. From SEM and X-ray diffraction, [1999Abe2] found a reversible transformation between 2 and

2´, which was confirmed by [2000Abe, 2000Tsa]. [1998Tak] indicates a considerable resemblance

between the hexagonal phase and the Laves phases MgZn2 leading to a quasi-identical parameter c. In the

hexagonal phase, two dimensional arrangement of clusters consist on columnar cage formed by Y atoms in

which the Zn and Mg atoms are tetrahedrally packed [1998Tak, 1999Abe1]. The icosahedral phase appears

to be realized as a results of a rearrangement of the column structure. According to [1998Sin], a two-fold

axis of the icosahedral phase corresponds to the six-fold axis of the hexagonal phase. The crystal lattice

parameters of the hexagonal phase 2, deduced from single crystal analysis [1998Tak], are reported in

Table 1. The crystal structure parameters of 2´ and 2´´ are also reported. The phase 2´ is identical to Z

detected by [1982Pad] and its composition of YMg3Zn6 was confirmed by [1997Tsa]. From electron

diffraction pattern, [2000Kou] concluded that an ordering of the icosahedral structure occurs in the

Y10-xMg30+xZn60, being face centred for x = 0 to simple structure for x = 4. The faced centred icosahedral

structure is considered as an order superstructure of the simple icosahedral atomic arrangement. This

confirms the observation done by [1999Abe1, 1999Abe2] of different types of quasi-crystal superstructures

for 2. The growth morphology of these phases is dodecahedral with visible pentagonal faces [1994Nii,

1998Fis]. [1999Abe2] suggested the existence of another hexagonal ternary phase Y7Mg27Zn66. No lattice

parameters for this phase were determined by [1999Abe2]. Most likely, it coincides with hexagonal phase,

established later by [2000Tsa] and designated by 3 in this assessment. The structure of the ternary phase

4 (Y2Mg3Zn3) has been determined by [1982Pad]. It has the MnCu2Al structure type (a = 684.8 pm) which

is in agreement with [1979Pad].

A phase Y(Mg, Zn)5 was also observed by [1997Tsa] as an intermediate phase for 2´ in the crystallization

process. Its crystallographic parameters reported by [1997Tsa] are presented in Table 1. Nevertheless,

[2000Tsa] showed the phase 3 as intermediate phase in the crystallization process for 2´ formation.

368


MSIT®

Mg–Y–Zn

For ternary phases some homogeneity ranges were shown, but their limits are not determined exactly.

From a Y2Mg97Zn1 alloy produced by extrusion of atomized powders at 300°C, [2001Ino] observed a novel

Mg based solid solution with a long periodic layered packing. The lattice parameters are a = 322 and

c = 3 521 pm.


[1982Pad, 1979Pad] state that a pseudobinary section exists between (Mg) and the ternary phase 1, and

contains a eutectic reaction at temperature of 540°C and composition of 24.2 mass% Y and 11.5 mass% Zn.

This has been accepted in the evaluation after shifting the assumed composition of 1.

An additional pseudobinary section Mg - 4 may exist according to the results of [1979Pad, 1982Pad].

From annealed sample at 300°C for 50 h, [2001Yi] indicates the existence of a pseudobinary eutectic

between (Mg) and 2´ in the Mg rich part (71 to 75 at.%) of the ternary Mg-Y-Zn system with Y varying

from 1 to 4 at.%.


The reaction scheme is presented in Fig. 1 and the ternary invariant equilibria are reported in Table 2.

[1979Pad] observed four four-phase invariant equilibria, three of them being eutectic and eutectoid

reactions, E1, E2, E4, and one being a transition reaction U1. In addition, the liquid undergoes two eutectic

decompositions L (Mg)+ 1 at 540°C (e2(max) reaction, Fig. 1) and L (Mg)+ 4 at 530°C (e3 reaction,

Fig. 1). The binary system Mg-Zn calculated by [1992Aga] presents three reactions p1, e4 and e5 as

described previously. The peritectic reaction p1 at 341.1°C and the eutectic reaction e4 at 341.0°C are

connected with a transition reaction U2 at ~340°C and a eutectic reaction E3 at ~340°C, which are presented

following the phase rule as two degenerate reactions D1 and D2 at ~340°C. Both reactions can not be

separated and, therefore, are shown together. For this reason these two four-phase reactions, U2 and E3

could not be observed by [1979Pad].

In the region Y > 7 at.% and Zn > 60 at.%, [2000Tsa] suggested that the 2 phase is formed via a peritectic

reaction which could be L+ 3 2 (at 598.4°C) or L+Y2Zn17 2. Moreover 2´ could be formed via the

peritectic reaction L+ 3 2´ (at 594.4°C). This last reaction is not in agreement with those reported by

[1997Tsa]: L+Y(Mg,Zn)5 2´ at 547°C with a liquid composition close to Mg7Zn3. [1997Lan1] also

reports a peritectic decomposition of 2´ at 600°C without specifying the reaction products.

The compositions of the liquid phase in the invariant reactions as far as found by [1982Pad] are given in

Table 2.

Liquidus Surface

The partial liquidus surface presented in Fig. 2 is based on the work of [1979Pad]. The liquidus isotherms

were constructed from the vertical sections given by [1979Pad, 1982Pad, 1977Dri] and the binary phase

diagrams given by [2001Shc, 2003Fab, Mas2]. Liquid isotherms and solid solubility isotherms of Y and Zn

in the Mg-rich region are also presented in a review done by [1977Ray] based on the experimental results

of [1968Zas]. However [1977Ray] states that the curves are of an unusual nature, and disagree with the data

of all other authors. So these isotherms and solid solubility curves have not been accepted here. The region

of primarily crystallization of 2´ as reported by [1997Lan2] is in agreement with this partial liquidus

surface.

A liquidus phase boundary (Mg)/ 2´ was located between the alloy compositions Y2Mg75Zn23 and

Yx(Mg74Zn26)100-x (x = 1-4 at.%) [2001Yi]. The same authors indicate the existence of two types of eutectic

microstructure on as-cast alloys depending on the previous thermal treatment.

Isothermal Sections

Figures 3 and 4 show the isothermal sections at 500 and 300°C, respectively, in the Mg-rich corner. These

isothermal sections have been adopted from [1979Dob] and have been extended according to the liquidus

surface and the vertical sections of [1979Pad] and [1977Dri]. There were minor discrepancies between

369


MSIT®

Mg–Y–Zn

those vertical sections and isothermal sections reported by [1979Dob]. The isothermal section at 300°C

presented by [1982Pad] gives additional equilibria between 2, 4, MgZn and MgZn2, but they are

inconsistent with the position of 2 and, therefore, could not be used.

Using annealed Zn-rich samples, [2000Tsa] constructed three isothermal sections at 427, 500 and 600°C

which involve quasicrystalline phases. According to the accepted binary phase diagrams, the phase fields

at the binary edges were adjusted and the Y2Zn17 phase was considered as stoichiometric. The isothermal

sections according to [2000Tsa] are reported in Figs. 5 to 7. The icosahedral 2´ phase region is close to

Y10Mg25Zn65 at 600°C and shifts to Y10Mg30Zn60 at 500 and 427°C. The composition of the 2´ phase at

600°C is identical to that of the 2 phase at 500°C. This confirms the existence of a phase transformation

between 2 and 2´ observed by [1999Abe2]. This phase transformation occurs between 600 and 500°C for

Y10Mg25Zn65. At 427°C, the liquid with a composition of Mg7Zn3 largely coexists with (Mg) and 2.

[2000Tsa] found also a W-phase with a composition range on the isothermal sections far from the nominal

composition reported by the same authors (Y3Mg2Zn3). As no other workers indicate the existence of this

phase, the phase fields with the W phase are not reported on the isothermal sections accepted in this

assessment.


Several workers investigated the mechanical properties on polycrystalline and single crystals of icosahedral

phases. Over the temperature range of 4 - 300 K the resistivity measured on the polycrystalline samples

[1995Kon] is always higher than that obtained from single crystals [1998Fis, 1999Fis]. In both cases, the

resistivity of Mg-Y-Zn quasicrystals has a very weak temperature dependence, with a small positive

magnetoresistance. Good elongation with relatively high strength is observed at room temperature

[2002Bae]. Others properties such as an ability to undergo plastic deformation [2000Yos], hardness

[2000Yos], magnetism [1999Kas, 2000Sch], electrical resistivity [1999Kas], thermal conductivity

[2000Gia] and tensile yield strength [2001Ino, 2002Pin] were also investigated in Mg-Y-Zn quasicrystals.

Mg base alloys with yttrium and zinc show high strength properties at common and elevated temperatures

[1980Rok].

Miscellaneous

[2000Ina] measured the heat capacity in the temperature range 4 - 300 K for an icosahedral phase of an

Y10Mg30Zn55 alloy. [2000Fis] studied growth of large single grains quasicrystals in the Mg-Y-Zn alloys.

[2000Est, 2000Let] used the icosahedral Mg-Y-Zn crystals for development of improved Brag scattering

technique in X-ray investigations. [2002Abe] found out the long-period structure in a high-strength

nanocrystalline 2Y-Mg-1Zn (at.%) alloy. [2002Kra] studied local atomic structure in Y8Mg42Zn50 alloy.

[2002Osh] investigated electronic structures of large approximants in Mg-Y-Zn alloys close the icosahedral

phase formation.

In as-cast and annealed Y5Mg50Zn45 alloy (300°C for 48 h) 2´ can also coexist with MgZn2 and MgZn

[1994Nii]. This was confirmed by [2000Tsa].

References

[1933Tar] Tarschisen, L., “X-Ray Investigation of the MgZn and MgZn5 Compounds”, Z. Kristallogr.,

A186, 423-438 (1933) (Crys. Structure, 6)


Magnesium and its Alloys”, Pergamon Press, London, New York, Paris, Los Angeles,


[1968Zas] Zaselyan, B.N., Saldau, P.Y., Afanasyev, S.K., “The Magnesium Rich Corner of the

Mg-Y-Zn Equilibrium Diagram”, Russ. Metall., 6, 130-133 (1968), translated from Izv.

Akad. Nauk. SSSR, Met., 6, 191, (1968) (Experimental, Equi. Diagram, 8)

370


MSIT®

Mg–Y–Zn

[1974Dri] Drits, M.E., Padezhnova, E.M., Miklina, N.V., “Phase Equilibria in Mg-Nd-Y-Zn Alloys”,

Russ. Metall., 4, 155-158 (1974), translated from Izv. Akad. Nauk. SSSR, Met., 4, 218-222,


[1977Dri] Drits, M.E., Padezhnova, E.M., Dobatkina, T.V., “Physicochemical Interactions of

Elements in Magnesium Alloys of the Mg-Y-M Systems (M=Mg, Nd, Al, Sid, Cd, Zn)”, in

“Magneviye Splavy”, Mater. Vses. Soveshch. Issled., Razrab. Primen. Magnievykh Plavov,

74-78, (1977) (Experimental, Equi. Diagram, 4)

[1977Ray] Raynor, G.V., “Constitution of Ternary and Some More Complex Alloys of Mg”, Int.

Metall. Rev., 22, 65-96 (1977) (Review, 83)

[1979Dob] Dobatkina, T.V., “Solid Solubility of Yttrium and Zinc in Magnesium”, Russ. Metall., 2,

169-173 (1979) (Experimental, Equi. Diagram, #, *, 9)

[1979Pad] Padezhnova, E.M., Melnik, E.V., Dobatnika, T.V., “Examination of Phase Equilibrium in

the Mg-Y-Zn System”, Russ. Metall., 1, 4179-182 (1979) (Experimental, Equi. Diagram,

Crys. Structure, #, *, 5)

[1980Rok] Rokhlin L.L., “Magnesium Alloys Containing Rare-Earth Metals” (in Russian), Nauka,

Moscow (1980) (Equi. Diagram, Crys. Structure, Review)



[1982Pad] Padeznhnova, E.M., Melnik, E.V., Miliyevskiy, R.A., Dobatkina, T.V., Kinzhibalo, V.V.,

“Investigation of the Mg-Zn-Y System”, Russ. Metall., 9, 185-188 (1982) (Experimental,

Equi. Diagram., Crys. Structure, #, *, 9)




[1992Luk] Lukas, H.L., Fries, S.G., “Demonstration of the Use of BINGSS with the Mg-Zn System”,

J. Phase Equilib., 13(5), 532-541 (1992) (Equi. Diagram, Thermodyn., 14)

[1993Luo] Luo, Z., Zhang, S., Tang, Y. Zhao, D., “Quasicrystals in As-Cast Mg-Zn-RE Alloys”, Scr.

Metall., 28, 1513-1518 (1993) (Crys. Structure, 12)

[1994Goe] Goedecke, T., Sommer, F., “Stable and Metastable Phase Equilibria in Mg-Zn2-Zn and



[1994Nii] Niikura, A., Tsai, A.P., Inoue, A., Masumoto, T., “Stable Zn-Mg-Rare-Earth Face-Centered

Icosahedral Alloys with Pentagonal Dodecahedral Solidification Structure”, Philos. Mag.

Lett., 69(6), 351-355 (1994) (Experimental, Crys. Structure, 9)

[1994Tsa] Tsai, A.P., Niikura, A., Inoue, A., Masumoto, T., “Highly Ordered Structure of Icosahedral

Quasicrystals in Zn-Mg-RE (RE=Rare Earth Metals) Systems”, Philos. Mag. Lett., 70(3),


[1995Kon] Kondo, R., Honda, Y., Hashimoto, T., Edagawa, K., Takeuchi, S., “Electrical Properties of

Zn-Mg-Y Icosahedral Quasicrystal”, Proc. Internat. Conf. on Quasicrystals, Avignon,


[1995Tsa] Tsai, A.P., Niikura, A., Inoue, A., Masumoto, T., “Icosahedral Quasicrystals in Zn-Mg-RE

(RE=Rare Earth Metal) System”, Proc. Internat. Conf. on Quasicrystals, Avignon, 628-635


[1997Fla] Flandorfer, H., Giovannini, M., Saccone, A., Rogl, P., Ferro, R., “The Ce-Mg-Y System”,

Metall. Mater. Trans. A, 28, 265-276 (1997) (Experimental, Equi. Diagram, Crys.

Structure, 24)

[1997Lan1] Langsdorf, A., Seuring, C. Ritter, F., Assmus, W., “Zn-Mg-(Y,RE) Quasicrystals: Growth

Experiments and Phase Diagrams”, Cryst. Res. Technol., 32(8), 1067-1072 (1997)


[1997Lan2] Langsdorf, A., Ritter, F., Assmus, W., “Determination of the Primary Solidification Area of

the Icosahedral Phase in the Ternary Phase Diagram of Zn-Mg-Y”, Philos. Mag. Lett.,

75(6), 381-387 (1997) (Experimental, Equi. Diagram, 9)

371


MSIT®

Mg–Y–Zn

[1997Tsa] Tsai, A.P., Niikura, A., Inoue, A., Masumoto, T., “Stoichiometric Icosahedral Phase in the

Zn-Mg-Y System”, J. Mater. Res., 12(6), 1468-1474 (1997) (Equi. Diagram,

Experimental, 6)

[1998Fis] Fisher, I.R., Islam, Z., Panchula, A.F., Cheon, K.O., Kramer, M.J., Canfield, P.C., Goldman,

A.I., “Growth of Large-Grain R-Mg-Zn Quasicrystals from the Ternary Melt (R = Y, Er, Ho

and Tb)”, Philos. Mag. B., 77(6), 1601-1615 (1998) (Experimental, Crys. Structure, 21)

[1998Lan] Langsdorf, A., Assmus, W., “Growth of Large Single Grains of the Icosahedral Quasicrystal

ZnMgY”, J. Cryst. Growth, 192, 152-156 (1998) (Experimental, Equi. Diagram, 18)

[1998Luk] Lukas, H.L., “Magnesium-Yttrium” in “COST-507: Thermochemical Database for Light

Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), European Communities,

Luxembourg, Vol.2, 224-226 (1998) (Thermodyn., 0)

[1998Sat] Sato, T.J., Abe, E., Tsai, A.P., “Composition and Stability of Decagonal Quasicrystals in the

Zn-Mg-Rare-Earth Systems”, Phil. Mag. Lett., 77(4), 213-219 (1998) (Experimental, 11)

[1998Sin] Singh, A., Abe, E., Tsai, A.P., “A Hexagonal Phase Related to Quasicrystalline Phases in

Zn-Mg-Rare-Earth System”, Phil. Mag. Lett., 77(2), 95-103 (1998) (Experimental, Crys.

Structure, 12)

[1998Tak] Takakura, H., Sato, A., Yamamoto, A., Tsai, A.P., “Crystal Structure of a Hexagonal Phase

and its Relation to a Quasicrystalline Phase in Zn-Mg-Y Alloys”, Philos. Mag. Lett., 78(3),


[1999Abe1] Abe, E., Takakura, H., Singh, A., Tsai, A.P., “Hexagonal Superstructures in the

Zn-Mg-Rare-Earth Alloys”, J. Alloys Compd., 283, 169-172 (1999) (Experimental, Crys.

Structure, 10)

[1999Abe2] Abe, E., Tsai, A.P., “Quasicrystal-Crystal Transformation in Zn-Mg-Rare-Earth Alloys”,

Phys. Rev. Let., 83(4) 753-756 (1999) (Experimental, Crys. Structure, 21)

[1999Fis] Fisher, I.R., Cheon, K.O., Panchula, A.F., Canfield, P.C., “Magnetic and Transport

Properties of Single-Grain R-Mg-Zn Icosahedral Crystals [R = Y, (Y1-xGdx), (Y1-xTbx), Tb,

Dy, Ho and Er]”, Phys. Rev. B, 59(1), 308-321 (1999) (Experimental, Phys. Prop., 30)

[1999Kas] Kashimoto, S., Matsuo, S., Nakano, H., Shimizu, T., Ishimasa, T., “Magnetic and Electrical

Properties of a Stable Zn-Mg-Ho Icosahedral Quasicrystal”, Solid State Commun., 109,


[2000Abe] Abe, E., Sato, T.J., Tsai, A.P., “Structure and Phase Transformation of the

Zn-Mg-Rare-Earth Quasicrystals”, Mater. Science Eng., 294-296, 29-32 (2000)


[2000Est] Estermann, M.A., Haibach, T., Steurer, W., Landsdorf, A., Wahl, M., “Weak Bragg

Scattering in Icosahedral Mg-Y-Zn”, Mater. Sci. Eng. A, 294-296, 237-241 (2000) (Crys.


[2000Fis] Fisher, I.R., Kramer, N.J., Islam, Z., Wiener, T.A., Kracher, A., Ross, A.R., Lograsso, T.A.,

Goldman, A.I., Canfield, P.C., “Growth of Large Single-Grain Quasicrystals from

High-Temperature Metallic Solutions”, Mater. Sci. Eng., A, 294-296, 10-16 (2000) (Crys.


[2000Gia] Gianno, K., Sologubenko, A.V., Chernikov, M.A., Ott, H.R., Fisher, I.R., Canfield, P.C.,

“Electrical Resistivity, Thermopower, and Thermal Conductivity of Single Grained (Y, Tb,

Ho, Er)-Mg-Zn Icosahedral Quasicrystals”, Mater. Sci. Eng., 294-296, 715-718 (2000)


[2000Ina] Inaba, A., Takakura, H., Tsai, A.P., Fisher, I.R., Canfield, P.C., “Heat Capacities of

Icosahedral and Hexagonal Phases of Zn-Mg-Y System”, Mater. Sci. Eng., 294-296,

723-726 (2000) (Experimental, 19)

[2000Kou] Kounis, A., Miehe, G., Fuess, H., “Investigation of Icosahedral Phases in the Zn-Mg-(Y,Er)

System by High Resolution Transmission Electron Microscopy”, Mater. Sci. Eng., 294-296,


[2000Let] Letoublon, A., Fisher, I,R., Sato, T.J., Boissieu, M., Boudard, M., Agliozzo, S., Mancini, L.,

Gastaldi, J., Canfield, P.C., Goldman, A.I., Tsai, A.-P., “Phason Strain and Structural

372


MSIT®

Mg–Y–Zn

Perfection in the Zn-Mg-Rare-Earth Icosahedral Phases”, Mater. Sci. Eng. A, 294-296,

127-130 (2000) (Crys. Structure, Experimental, Phys. Prop.,15)

[2000Luo] Luo, Z.P., Zhang, S.Q., “High-Resolution Electron Microscopy on the X-Mg12ZnY Phase

in High Strength Mg-Zn-Zr-Y Magnesium Alloy”, J. Mater. Sci. Let., 19, 813-815 (2000)

(Experimental, Crys. Structure)

[2000Sch] Scheffer, M., Rouijaa, M., Suck, J.B., Sterzel, R., Lechner, R.E., “Magnetic Neutrron

Scattering from Quasicrystalline Zn-Mg-Ho and Zn-Mg-Y at Low Temperatures”, Mater.

Sci. Eng., 294-296, 488-491 (2000) (Experimental, Magn. Prop., 12)

[2000Ste] Sterzel, R., Dahlmann, E., Langsdorf, A., Assmus, W., “Preparation of Zn-Mg-Rare Earth

Quasicrystals and Related Crystalline Phases”, Mater. Sci. Eng., 294-296, 124-126 (2000)


[2000Tsa] Tsai, A.P., Murakami, Y., Niikuras, A., “The Zn-Mg-Y Phase Diagram Involving

Quasicrystals”, Philos. Mag. A, 80(5), 1043-1054 (2000) (Experimental, Equi.

Diagram, 17)

[2000Yos] Yoshida, T., Itoh, K., Tamura, R., Takeuchi, S., “Plastic Deformation and Hardness in

Mg-Zn-(Y, Ho) Icosahedral Quasicrystals”, Mater. Sci. Eng., 294-296, 786-789 (2000)


[2001Ino] Inoue, A., Kawamura, Y., Matsushita, M., Hayashi, K., Koike, J., “Novel Hexagonal

Structure and Ultrahigh Strength of Magnesium Solid Solution in the Mg-Y-Zn System”,

J. Mater. Res., 16(7), 1894-1900 (2001) (Experimental, Mechan. Prop., 8)

[2001Sat] Sato, T.J., Abe, E., Tsai, A.P., “Decogonal Quasicrystals in the Zn-Mg-R alloys (R =

rare-earth and Y)”, Mater. Sci. Eng. A, 304-306, 867-870 (2001) (Experimental, Crys.

Structure, 9)

[2001Shc] Shcherban, O., Ilyenko, S., “Mg-Zn (Magnesium-Zinc)” in “Ternary Alloys: A

Comprehendium of Evaluated Consitutional Data and Phase Diagrams”, Effenberg, G.,

Aldinger, F., Rogl, P. (Eds.), MSI GmbH, Stuttgart, (2001) (Crys. Structure, Equi. Diagram,

Assessment, 9)

[2001Yi] Yi, S., Park, E.S., Ok, J.B., Kim, W.T., Kim, D.H., “(Icosahedral Phase + Mg) Two Phase

Microstructures in the Mg-Zn-Y Ternary System”, Mat. Sci. Eng., A, 300, 312-315 (2001)


[2002Abe] Abe, E., Kawamura, Y., Hayashi, K., Inoue, A., “Long-Period Ordered Structure on a a

High-Strength Nano-Crystalline Mg-1at%Zn-2at%Y Alloy Studied by Atomic-Resolution

Z-Contrast STEM”, Acta Mater., 50, 3845-3857 (2002) (Crys. Structure, Experimental, 20)

[2002Bae] Bae, D.H., Kim, S.H., Kim, D.H., Kim, W.T., “Deformation Behavior of Mg-Zn-Y Alloys

Reinforced by Icosahedral Quasicrystalline Particles”, Acta Mater., 50, 2343-2356 (2002)


[2002Kra] Kramer, M.J., Hong, S.T., Canfield, P.C., Fisher, I.R., Corbett, J.D., Zhu, Y., Goldman, A.I.,

“The Local Atomic Structure of R-Mg-Zn (R = Y, Gd, Dy, and Tb)”, J. Alloys Compd.,

342(1-2), 82-86 (2002) (Crys. Structure, Experimental, 14)

[2002Osh] Oshio, K., Ishii, Y., “Electronic Structures of Large Approximants of Zn-Mg-Y”, J. Alloys

Compd., 342, 402-404 (2002) (Calculation, Crys. Structure, 9)

[2002Pin] Ping, D.H., Hono, K., Kawamura, Y., Inoue, A., “Local Chemistry of a Nanocrystalline

High-Strength Mg97Y2Zn1 Alloy”, Philos. Mag. Lett., 82(10), 543-551 (2002)


[2003Fab] Fabrichnaya, O., Pisch, A., “Mg-Y (Magnesium-Yttrium)”, MSIT Binary Evaluation

Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International

Services GmbH, Stuttgart; Document ID: 20.15845.1.20 (2003) (Equi. Diagram, Crys.

Structure, Assessment, 20)

373


MSIT®

Mg–Y–Zn


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.944

c = 521.076

a = 320.99 to

319.57

c = 521.08 to

518.82

pure Mg at 25°C [V-C2, Mas2]

dissolves up to 3.75 at.% Y at 566°C,

up to 2.4 at.% Zn at 340°C [Mas2]

at 0 to 2.81 at.% Zn [V-C2]

( Y)(h)

1522 - 775

cI2

Im3m

W

a = 407 lattice parameter estimated [Mas2]

( Y)(r)

< 1478

hP2

P63/mmc

Mg

a = 364.82

c = 573.18

at 25°C [Mas2]

(Zn)

< 419.5

hP2

P63/mmc

Mg

a = 266.47

c = 494.69



Mg51Zn20

341.1 - 325

oI158

Immm

Mg51Zn20

a = 1408.3 0.3

b = 1448.6 0.3

c = 1402.5 0.3

at 28.17 at.% Zn [1992Aga, 1992Luk]

[1981Hig, V-C2]

labeled as Mg7Zn3 in [Mas2, 1992Aga,

1992Luk]

MgZn

< 347

oP48

?

a = 533

b = 923

c = 1716



[1959Ray]



MgZn2

< 586

hP12

P63/mmc

MgZn2 a = 522.3 0.1

c = 856.6 0.3

at 66-67 at.% Zn at 381°C [1994Goe]

[V-C2]

Y5Mg24+x

< 616.6

cI58

I43m

Mn

a = 1127.8 to

1125.0

84 to 87 at.% Mg [V-C2]

~14.24 to 17.24 at.% Y [Mas2]

YMg2

< 780

hP12

P63/mmc

MgZn2

a = 603.7 0.1

c = 975.2 0.2

[V-C2]

Y2Zn17

< 860

hR19

R3m

Th2Zn17

a = 897.19 0.05

c = 1314.14 0.08

at 89.5 at.% Zn

[V-C2]

374


MSIT®

Mg–Y–Zn

YMgxZn1-x

YMg

< 935

YZn

< 1105

cP2

Pm3m

CsCl

a = 362

a = 381.0 to 378.1

a = 356

0 < x < 1

unknown composition [1982Pad]

~50 to 52 at.% Mg [V-C2]

[V-C2]

* 1, ~Y9Mg85Zn6 or

YMg12Y

h*

or

trigonal

a = 322.4

c = 4698.5

a = 1577.2

= 11.73°

[2000Luo]

* 2,

~Y6.86Mg27.92Zn65.22

h*

P63/mmc a = 1457.9 0.2

c = 868.7 0.1

[1999Abe2]

92 atoms [1998Tak]

* 2´, ~YMg3Zn6 icosahedral

Fm53

a = 519

quasicrystalline phase

small homogeneity range [1993Luo]

[2000Tsa]

* 2´´, ~Y2Mg38Zn60 decagonal

P105/mmc

a = 227

c = 255

quasicrystalline phase

Composition close to Mg2Zn3 [1998Sat]

* 3, ~Y15Mg15Zn70 h*

P63/mmc

a = 776

c = 920

[2000Tsa]

* 4, ~Y2Mg3Zn3 cF16

Fm3m

MnCu2Al a = 684.8

[1982Pad]

homogeneity range exists

[V-C2]

* Y7Mg27Zn66 h* - [1999Abe2]

* Y(Mg,Zn)5 m* a = 850

b = 650

c = 950

= 67.2°

[1997Tsa], doubtful

W, Y3Mg2Zn3 c*

Fm3m

a = 683 [2000Tsa]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

375


MSIT®

Mg–Y–Zn



Mg Y Zn

L, (Mg), MgZn, Mg51Zn20, 2´ 340 D1 L

(Mg)

MgZn

Mg51Zn20

2´

70.6

100.0

50.0

71.8

30.0

0.3

0.0

0.0

0.0

10.0

29.1

0.0

50.0

28.2

60.0

L (Mg) + 1 + 4 527 E2 L

(Mg)

1

4

82.3

100.0

85.0

37.5

7.2

0.0

9.0

25.0

10.5

0.0

6.0

37.5

L (Mg) + 4 530 e3 (max) L

(Mg)

4

82.3

100.0

37.5

7

0.0

25.0

10.7

0.0

37.5

L (Mg) + Y5Mg24+x + 1 533 E1 L

(Mg)

Y5Mg24+x (x = 0)

1

87.7

100.0

82.8

85.0

9.2

0.0

17.2

9.0

3.1

0.0

0.0

6.0

L (Mg) + 1 540 e2 (max) L

(Mg)

1

85.5

100.0

85.0

8.8

0.0

9.0

5.7

0.0

6.0

Fig. 1: Mg-Y-Zn. Partial reaction scheme

Mg-Y A-B-CMg-Y-Zn Mg-Zn

l (Mg)+Y5Mg

24+x

566 e1

L (Mg) + τ1

540 e2(max)

L (Mg)+Y5Mg

24+x+τ

1533 E

1

L (Mg) + τ4

530 e3(max)

L (Mg) + τ1

+ τ4

527 E2

L + τ4

(Mg) + τ´2

448 U1

l+(Mg) Mg51

Zn20

341.1 p1

l Mg51

Zn20

+MgZn

341 e4

Mg51

Zn20

(Mg)+MgZn

325 e5

L ,(Mg),MgZn,Mg51

Zn20

,τ´2

~340 D1,D

2

Mg51

Zn20

MgZn+(Mg)+τ´2

325 E3

(Mg)+Y5Mg

24+x+τ

1 L+τ1+τ

4

?

(Mg) + τ1

+ τ4

MgZn+(Mg)+τ´2

?

?

(Mg)+τ´2+τ

4(Mg)+L+τ´

2

MgZn+Mg51

Zn20

+τ´2(Mg)+Mg

51Zn

20+τ´

2

L+τ´2+MgZn

L+Y5Mg

24+x+τ

1

L+τ´2+τ

4

376


MSIT®

Mg–Y–Zn

10

20

30

40

60 70 80 90

10

20

30

40

Y 50.00

Mg 50.00

Zn 0.00

Mg

Y 0.00

Mg 50.00

Zn 50.00Data / Grid: at.%

Axes: at.%

Y5Mg

24+x

E1

E2

U1

MgZn

p1,e

4, 341°C

600

550

500

450

400

τ4

U2, E

3

(341.1>T>340°C)

e1

e3

e2

(Mg)

10

20

80 90

10

20

Y 30.00

Mg 70.00

Zn 0.00

Mg

Y 0.00

Mg 70.00


Axes: at.%

(Mg)

L + (M

g)

LL + τ4

τ1

L + (Mg) + τ

4

(Mg) + τ

1 + τ4

(Mg) + τ1

Y5Mg

24+x

(Mg) + Y5Mg

24+x + τ

1

Fig. 2: Mg-Y-Zn.

Partial liquidus

surface

Fig. 3: Mg-Y-Zn.

Partial isothermal

section at 500°C

377


MSIT®

Mg–Y–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Y Mg


Axes: at.%

Y2Zn

17

τ2

τ2'

τ3

L

20

40

60

40 60 80

20

40

60

Y 80.00

Mg 20.00

Zn 0.00

Mg

Y 0.00

Mg 20.00


Axes: at.%

Y5Mg

24+x

τ1

τ4

τ2

MgZn

(Mg)

(Mg)+τ

2 +τ4

(Mg)+

MgZ

n+τ2

Fig. 5: Mg-Y-Zn.


600°C

Fig. 4: Mg-Y-Zn.

Partial isothermal

section at 300°C

378


MSIT®

Mg–Y–Zn

20

40

60

80

20 40 60 80

20

40

60

80

Y Mg


Axes: at.%

Y2Zn

17

τ3

τ2'

τ2

MgZn2

L

20

40

60

80

20 40 60 80

20

40

60

80

Y Mg


Axes: at.%

Y2Zn

17

MgZn2

τ3

τ2'

τ2

L

(Mg)

Fig. 6: Mg-Y-Zn.


500°C

Fig. 7: Mg-Y-Zn.


427°C

379


MSIT®

Mg–Zn–Zr

Magnesium – Zinc – Zirconium

Hans Jürgen Seifert

Literature Data

Only the Mg-rich corner of the Mg-Zn-Zr system was investigated in literature covering a Zr-content of up

to 4 mass% and a Zn-content of up to 10 mass%. Most of the publications investigate the influence of Zn

on the grain refinement of Zr on Mg.

[1957Ich, 1959Ich1, 1959Ich2] detected the solubility of Zr in liquid Mg under the influence of Zn using

different experimental techniques. [1957Ich] added 5 and 10 mass% Zr, respectively, to liquid Mg

(>99.99%). Zr was added as ZrCl4. 1, 2 or 4 mass% Zn was added. The Mg was molten in MgO-covered

steel crucibles using a flux. Zn was added first at 700°C and subsequently Zr was added at 725 or 825°C,

respectively. The melts were kept 30 min at 700 or 800°C and were subsequently water-cooled. The samples

were analyzed by metallography and chemical analysis. [1959Ich1, 1959Ich2] melted very high purity Mg

and added pure Zn (1-6 mass%) and subsequently Mg-20%Zr alloys (total amount of Zr was up to 3

mass%). The melts were kept 30 minutes at 700 or 800°C, respectively and water-quenched. The samples

were investigated by chemical analysis, metallography and X-ray. Results of [1957Ich, 1959Ich1,

1959Ich2] show, that Zn was scarcely effected to the solubility of Zr in liquid Mg. The solubility of Zr at

700°C was 0.75-0.8 mass% and about 0.9 mass% at 800°C. X-ray analysis of the samples gave small

impurities of Fe2Zr, ZrC, ZrH, and ZnO. Mainly uncompounded metallic Zr was found when alloying Zr

over the solubility limit in molten Mg-Zn-Zr alloys. [1968Bab] derived Mg-Zr alloys by adding a

Mg+15%Zr alloy (10 mass%) to liquid Mg at 780°C. Zn was added as pure metal at the same temperature.

No information is given on the purity of the samples. The alloys were melted in iron crucibles using a flux

and held at 780-800°C for 15-20 minutes before pouring in open cast iron moulds heated to 150-200°C and

subsequently cooled. The authors did not find a significant effect of Zn on the solubility in liquid Mg and

found a maximum value of 0.5 mass% Zr.

Ternary alloys containing between 1 and 44 mass% Zn and 0.6-1.1 mass% Zr were prepared by [1969Las]

using a flux and iron crucibles. The starting materials were magnesium grade MG (GOST 804-56), zinc

grade U1 (GOST 3640-47), and a master alloy Mg-20 mass% Zr. The zirconium was introduced in all the

alloys in an amount equivalent to 2% of the charge weight. The specimen were cast in sand moulds. As-cast

samples (1) were investigated and specimen after aging for 6 h at 300°C (2), after heating at 440°C followed

by cooling in air (3), and after heating at 440°C for 6 h cooling in air and aging at 150°C for 50 h (4). The

chemical and phase compositions of all Mg-Mn-Zr alloys are presented. In the heat treated samples the

authors found Mg, Mg7Zn3, Mg2Zn3, MgZn, Zr3Zn2, ZrZn. Mg7Zn3 was reported to be stable at high

temperatures only, MgZn is stable between 300°C and the temperature at which the Mg7Zn3 phase starts to

be stable. Results of [1969Las] do not represent stable phase equilibria. [1975Loe] used Mg, (> 99.9%),

with exactly analyzed impurities (Si, Mn, Al, Fe, Ni and Cu) and Zink of 99.99% purity. Zr was introduced

by an alloy of 35% Zr and 55% Mg and reaction salt. The alloys were molten in Fe-crucibles. After melting

of salt covered Mg the temperature was increased to 800°C and the pre-alloys and elements were added. Zr

was added as last element. Liquidus points were fixed by slow cooling the melt in the crucible. The samples

were investigated by SEM/EDAX. The liquid was over-saturated with Zr at 850°C and subsequently cooled

down to 800°C and the sample analyzed after one hour holding time. The same procedure was repeated at

730 and 660°C. Afterwards, the sample was slowly cooled down and investigated by thermal analysis. Zr

was detected by photometric methods. From these data the liquidus surface in the Mg-rich corner of the

system could be constructed. According to the experimental results, the solubility of Zr in the liquid

increases up to 5 mass% Zn and decreases at higher Zn-content. The peritectic temperature in the Mg-Zr

system is decreased by adding Zn. The steep liquidus becomes more flat with more than 5% Zn. The shape

of the monovariant line was confirmed by Mg grain size measurements. [1978Dri] investigated the Mg-rich

corner of the system up to 10 mass% Zn and 2 mass% Zr. The purity of the initial elements was 99.975%

Mg, 99.96% Zn. No impurity data are given for Zr. The samples were prepared by warm pressing using a

380


MSIT®

Mg–Zn–Zr

flux and were subsequently encapsulated under Ar. They were heat treated at 500°C (50 h) or at 300°C

(500 h) and quenched in water. The samples were investigated by X-ray diffraction, metallography and

thermal analysis (heating and cooling). An isothermal section at 300°C was presented. The Zn and Zr

solubility in (Mg) in the three phase field (Mg)+(Zr)+Zn2Zr3 was detected to be 1.5 Zn, 0.1 Zr (mass%) at

300°C and 1.5 Zn, 0.3 Zr (mass%) at 500°C. An isopleth from 10Zn-0Zr to 2Zr-0Zn was determined.

[1978Mor] prepared alloys by “standard techniques”. No information is given on the experimental

procedure or the purity of the samples. The alloys were investigated by X-ray, chemical analysis and

metallography. As-cast alloys and after heat treatment alloys were investigated. The commercial alloys

were contaminated with hydrogen and/or Cd/Nd and cannot be considered for the assessment of the ternary

Mg-Zn-Zr. The experimental alloy was heat treated for 6 h at 300°C and subsequently air cooled or heat

treated for 6 h at 440°C, air cooled and subsequently aged at 150°C for 50 h. Phase composition of the

residues is documented and deviates from the results of [1978Dri]. Obviously the samples are not in

equilibrium. [1980Zak] assessed experimental results of [1969Las] and [1978Dri] and constructed the

liquidus surface of the complete Mg-Zn-Zr system. Additionally, solid solution data of Zn and Zr,

respectively, in the Mg-rich corner at 300°C are presented. [1993Bha] gave a review of the system accepting

the liquidus surface given by [1975Loe] and not discussing the results of [1978Dri].

Binary Systems

For the Mg-Zn binary system the diagram of [1992Aga, 1992Luk] was accepted taking into account later

refinement by [1994Goe] regarding the temperature of the eutectic reaction l Mg2Zn11+(Zn). The phase

diagrams for the binary systems Mg-Zr and Zn-Zr according to [Mas2] were accepted. However Zn2Zr3 was

considered to be a stable phase in the Zn-Zr system.

Solid Phases

The composition of Mg7Zn3 given by [1971Dri] was changed to the accepted composition Mg51Zn20,

which arises from X-ray investigations. No ternary phases have been found. All solid phases are listed in

Table 1.


[1978Dri] reported the following invariant equilibria:

1. L + Zr (Mg) + Zn2Zr3 (570°C);

2. L (Mg) + Zn2Zr3 + Mg7Zn3 (340°C);

3. Mg7Zn3 (Mg) + MgZn + Zn2Zr3 (335°C).

However, these data were not accepted, as they are not consistent with the binary phase diagrams and the

liquidus surface given by [1975Loe].

Liquidus Surface

The reported partial liquidus surface according to [1975Loe] was accepted and is shown in Fig. 1. The

liquidus of the isopleth, given by [1978Dri] is not consistent with the accepted binary systems Mg-Zn and

Mg-Zr systems and the liquidus surface of [1975Loe]. According to [1975Loe] the solubility of Zr in liquid

Mg is slightly increased by increasing Zn amount which is not in accordance with data of [1957Ich,

1959Ich1, 1959Ich2]. Results of [1975Loe] were accepted here, because the well documented experimental

procedure is very suitable for determining the data reported. The liquidus surface construction presented by

[1980Zak] is mainly based on assumptions on possible phase equilibria and is not taken into account in the

present assessment.

Isothermal Sections

Figure 2 shows the slightly redrawn isothermal section at 300°C according to [1978Dri]. This isothermal

section is in accordance with binary systems.

381


MSIT®

Mg–Zn–Zr

Miscellaneous

Only few equilibrium data on the Mg-Zn-Zr system are available. Most of the publications report

investigations on alloys including metastable phases with high amount of impurities and were not accepted.

The invariant equilibria, partly reported by [1978Dri] need to be re-investigated experimentally. There is

only a small influence of Zn on the solubility of Zr in liquid Mg and Zr can be used for grain refinement in

these alloys.

References

[1933Tar] Tarschisch, L., “X-Ray Investigation of the MgZn and MgZn5 Compounds” (in German),

Z. Kristallogr., 86, 423-438 (1933) (Crys. Structure, 6)

[1957Ich] Ichikawa, R., “The Solubility of Zirconium in Magnesium and its Alloys at Liquidus State,

1st Part: On Alloys Containing Aluminium and Zinc” (in Japanese), Bull. Nagoya Inst.

Technol., 9, 902-905 (1957) (Experimental, 7)

[1959Ich1] Ichikawa, R., “The Intermetallic Compound Formed by Impurities and Zr, 2nd Report: On

Al, Fe, Mn and Si as Impurities in Mg-Zr Alloys” (in Japanese), Nippon Kinzoku

Gokkai-Shi, 23, 192-194 (1959) (Experimental, 5)

[1959Ich2] Ichikawa, R., “Suitable Alloying Elements for Mg-Zr Alloys”, Nippon Kinzoku Gakkai-Shi,

23, 612-616 (1959) (Experimental, 10)

[1959Ray] Raynor, G.V., “The Physical Metallurgy of Mg and its Alloys”, London: Pergamon, 531pp.

(1959) (Review)

[1968Bab] Babkin, V.M., “Solubility of Zirconium in Liquid Magnesium and the ML5 Alloy”, Met.

Sci. Heat Treat., 221-223 (1968), translated from Metalloved. Term. Obrab. Met., 3, 61-64


[1969Las] Lashko, N.F., Morozova, G.I., Andreyeva, F.S., Tikhonova, V.V., Gerasimova, M.A.,

“Phase Composition of Cast Mg-Zn-Zr Alloys”, Russ. Metall., 129-132 (1969), translated

from Izv. Akad. Nauk SSSR, Met., 2, 159-163 (1969) (Experimental, 7).

[1971Dri] Drits, M.E., Padezhnova, E.M., Miklina, N.V., “The Phase Diagram of the Mg-Nd-Zn

System in the Magnesium-Rich Region” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn.

Metall., 4, 103-107 (1971) (Equi. Diagram, Experimental, 8)

[1975Loe] Löhberg, K., Schmidt, G., “Contribution to the Grain Refinement Effect of Zr in Mg Alloys”

(in German), Giessereiforschung, 27, 75-82 (1975) (Equi. Diagram, Experimental, #, *, 36)

[1975Yar] Yarmolyuk Y.P., Kripyakevich P.I., Mel’nik E.V., “Crystal Structure of the Compound

Mg4Zn7”, Sov. Phys. Crystallogr., 20(3), 329-331 (1975)

[1978Dri] Drits, M.E., Guzei, L.S., “The Magnesium-Zinc-Zirconium System in the Magnesium Rich

Region” (in Russian), Probl. Metalloved. Tsv. Splavov, 81-89 (1978) (Equi. Diagram,

Experimental, #, 7)

[1978Mor] Morozova, G.I., Tikhonova, V.V., Lashko, N.F., “Phase Composition and Mechanical

Properties of Cast Mg-Zn-Zr Alloys”, Met. Sci. Heat Treat., 657-600 (1978), translated

from Metalloved. Term. Obrab. Met., 52-54 (1978) (Experimental, 6)

[1980Zak] Zakharov, A.M., “The System Mg-Zn-Zr”, Promyshlennye Splavy Tsvetnykh Metallov” (in

Russian), in “Industrial Nonferrous Alloys: Phase Diagrams”, 101-103 (1980) (Review, 3)






[1992Luk] Lukas, H.L., Fries, S.G., “Demonstration of the Use of BINGSS with the Mg-Zn System”,

J. Phase Equilib., 13(5), 532-541 (1992) (Equi. Diagram, Thermodyn., 14)

[1993Bha] Bhan, S., Lal, A., “The Mg-Zn-Zr System”, J. Phase Equilib., 14, 634-637 (1993)

(Review, 21)

382


MSIT®

Mg–Zn–Zr





Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Mg)

< 650

hP2

P63/mmc

Mg

a = 320.944

c = 521.076

a = 320.99 to

319.57

c = 521.08 to

518.82


dissolves 2.4 at.% Zn at 340°C [Mas2] at

0 to 2.81 at.% Zn [V-C2]

( Zr)

< 1855

cI2

Im3m

W

a = 413 [V-C2]

( Zr)

< 863

hP2

P63/mmc

Mg

a = 365.66

c = 1179.83

[V-C2]

(Zn)

< 419.5

hP2

P63/mmc

Mg

a = 266.47

c = 494.69



Mg51Zn20

341.1 - 325

oI158

Immm

Mg51Zn20

a = 1408.3 0.3

b = 1448.6 0.3

c = 1402.5 0.3

at 28.17 at.% Zn [1992Aga, 1992Luk]

[1981Hig, V-C2]

labeled as Mg7Zn3 in [Mas2, 1992Aga,

1992Luk]

MgZn

< 347

oP48

?

a = 533

b = 923

c = 1716



[1959Ray]



Mg2Zn3

< 416

mC110

C2/m

Mg4Zn7 a = 2596

b = 524

c = 2678

= 148.6°



[1975Yar, V-C2]

labeled as Mg4Zn7 in [V-C2]

MgZn2

< 586

hP12

P63/mmc

MgZn2 a = 522.3 0.1

c = 856.6 0.3

at 66.0-67.0 at.% Zn at 381°C [1994Goe]

[1994Goe]

[V-C2]

Mg2Zn11

< 381

cP39

Pm3

Mg2Zn11 a = 855.2 0.5

at 84.1-84.6 at.% Zn at 368°C [1994Goe]

[Mas2, 1994Goe]

[V-C2]

383


MSIT®

Mg–Zn–Zr

ZnZr2 tI6

I4/mmm

MoSi2

a = 330.3

c = 1126.0

[V-C]

Zn2Zr3 tP20

P42nm

Al2Gd3

a = 763.3

c = 696.5

[V-C]

ZnZr

< 1110

cP2

Pm3m

CsCl

a = 333.6 [V-C]

Zn2Zr

< 1180

cF24

Fd3m

Cu2Mg

a = 739.4 [V-C]

´Zn3Zr

< 910

t* a = 816.0

c = 1623.0

[V-C]

´Zn3Zr

< 1100

c* - [Mas2]

Zn6Zr

< 750

t* - [Mas2]

Zn14Zr

< 545

cF184 - [Mas2]

Zn22Zr cF184

Fd3m

Al18Cr2Mg3

a = 1410.3 [V-C]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

384


MSIT®

Mg–Zn–Zr

1.61.20.80.4

MgZr

Mg

Zn

2.00

98.00

0.00

4.0

8.0

12.0

16.0

660

650

640

640

630

630

20.0

2.0

Zr, mass%

Zn,m

ass%

0.2 2.0

800

730

Fig. 1: Mg-Zn-Zr.

Liquidus surface

Mg Zr

Mg

Zn

4.00

96.00

0.00

0.8 1.6 2.4 3.2

4.0

8.0

12.0

16.0

(Mg)

(Mg) + Zr

(Mg) + Zr + Zn Zr2 3

(Mg) + Zn Zr2 3

(Mg) + + MgZnZn Zr2 3

Zr, mass%

0

Zn,m

ass%

4.0

20.0Fig. 2: Mg-Zn-Zr.


300°C

385


MSIT®

Mo–Si–Ti

Molybdenum – Silicon – Titanium

Anatoliy Bondar and Hans Leo Lukas

Literature Data

The Mo-Si-Ti system was studied at temperatures from 1250 or 1300°C to melting due to the fact that the

Mo-Si-Ti silicides are promising candidate materials for high-temperature structural applications as well as

for oxidation resistant electric heating elements.

[1952Now] studied the TiSi2-MoSi2 section at 1300°C, the melting behavior of the complex disilicides

[1956Kud] and the Ti5Si3-Mo5Si3 section [1954Sch]. Samples were prepared by hot pressing at 1400 to

1500°C (TiSi2-MoSi2) and 1500 to 1700°C (Ti5Si3-Mo5Si3) using powders of Si (99.7 mass% purity),

titanium hydride and metals as starting materials. The samples were annealed at 1300°C for 20 h

(disilicides) and at 2000°C for 2 h (M5Si3 silicides) in hydrogen atmosphere and examined by powder XRD,

metallography and pyrometric measurements of melting points.

[1965Gar] investigated the reaction of liquid Ti with molybdenum disilicide, MoSi2. They identified

equilibrium between a (Ti,Mo) solid solution, Ti5Si3 and a Mo rich silicide, either Mo3Si or Mo5Si3, in the

reaction zone after heating at 1320°C for 3 h.

[1971Koc1, 1972Sve1, 1972Sve2, 1974Sve] examined the ternary system in more detail. They prepared

227 ternary alloys in the Ti5Si3-Si-Mo5Si3 region and investigated them by light metallography (in selected

cases using SEM/EMPA), X-ray diffraction (Debye-Scherrer technique, V and Cu radiations) and DTA

with string W/80W-20Re (mass%) thermocouples [1971Koc1, 1971Koc2] (Al2O3, ZrO2, HfO2, or BeO

crucibles and heating rates of 30 to 50°C min-1 were used). Before this work they had studied in the same

manner 63 Mo-Si alloys and approximately 30 Si-Ti alloys [1970Sve1, 1970Sve2, 1971Sve1, 1971Sve2].

The alloys were arc-melted from iodide Ti of 99.86 mass% purity, bulk Mo of 99.9 mass% and single crystal

Si (of “KSD” grade). The TiSi2-Si-MoSi2 alloys were annealed in purified argon at 1250°C for 70 h (more

refractory alloys were kept 180 h in addition). The TiSi2-MoSi2 alloys were annealed at higher temperatures

(1725°C/15 h, 1810°C/10 h, 1875°C/5 h and 1920°C/2 h) or in two steps: after homogenization at 1500°C

for 58 h (more refractory alloys at 1600°C for 140 h), they were annealed at 1425°C for 150 h or at 1300°C

for 250 h. The Ti5Si3-TiSi2-MoSi2-Mo5Si3 alloys were annealed at 1425°C for 75 h (the alloys containing

more 40 mass% Mo were previously annealed at 1600°C for 25 h). The Ti5Si3-Mo5Si3 alloys were

homogenized at 1800°C for 25 h followed by annealing at 1425°C for 75 h. The obtained data were

presented by isothermal sections at 1425°C (Ti5Si3-TiSi2-MoSi2-Mo5Si3) and 1250°C (TiSi2-Si-MoSi2),

the TiSi2-MoSi2 isopleth, the liquidus and solidus surfaces of the TiSi2-Si-MoSi2 region and several vertical

sections for that region.

[1991Fra, 1992Boe, 1998Fra] studied TiSi2-MoSi2 alloys, as well as the binary Mo-Si system in the vicinity

of MoSi2. The alloys were arc-melted from high purity elements, Ti of 99.7 mass%, Mo of 99.95 mass%

and Si of 99.9999 mass%. The arc processed alloys were annealed at T > 1300°C and then examined by

powder X-ray diffraction and SEM/EMPA (EDS). The temperature of the reaction Lp+(MoSi2) (Ti,Mo)Si2was determined by heat treatment of samples in gettered He or Ar followed by quenching at ~400°C min-1

and further metallographic examination and EPMA estimation of the liquidus composition. The

Lp+(Ti,Mo)Si2 TiSi2 peritectic reaction was studied by DTA.

[2001Wei] melted TiSi2-MoSi2 alloys in an arc furnace (no data there were given on the purity of starting

materials) and homogenized them at 1400°C for 168 h. Melting points were measured by DTA running up

to 1650°C. The samples were studied by SEM/EPMA (EDS) and powder X-ray diffraction.

[2003Yan, 2004Yan] prepared 46 ternary alloys (from 10 to 63at.% Si) by non-consumable electrode arc

melting from high purity elements (99.995 mass% Ti, 99.95 % Mo and 99.9999 % Si). Samples were

annealed in ultra-pure argon at 1600°C for 150 h and at 1425°C for 300 h and examined by SEM/EMPA

(EDS) techniques and X-ray diffractometry. Using Thermo-Calc and Pandat softwares, [2003Yan]

optimized a complete thermodynamic description of the ternary system considering besides their own ones

also the earlier experimental results of [1972Sve2, 1991Fra, 1992Boe]. They adopted the thermodynamic

386


MSIT®

Mo–Si–Ti

descriptions of the binary systems from [1996Sei] (Si-Ti, version with Ti5Si3 simplified as stoichiometric

phase), [1998Sau] (Mo-Ti) and [2000Liu] (Mo-Si). The obtained dataset is compared with experimental

data by calculation of isothermal sections at 1425 and 1600°C, the liquidus projection and the TiSi2-MoSi2section. For several alloy compositions (Ti,Mo)3Si the solidification path was calculated after the Scheil

model and compared with micrographs of the corresponding as-cast alloys.

Binary Systems

The unary and binary phases are presented in Table 1. The Mo-Ti phase diagram is taken from [Mas2] after

Murray [1987Mur1].

A complete evaluation of the Si-Ti system was done by [1987Mur2] represented in [Mas2]. A

thermodynamic optimization was carried out by [1996Sei], which is applied to the Mo-Si-Ti system in

[2003Yan]. Good agreement is observed between the experimental data from literature and the calculated

phase diagram, with the exception of the TiSi2 melting point where the calculation is noticeably lower than

the experimental data. So, the difference between the TiSi2 congruent melting temperature and the TiSi+

TiSi2 eutectic is 1478-1474 (°C) in [1996Sei] and 1500-1470 (°C) in [1970Sve1, 1971Sve2]. More recent

data of [1998Du] 1488-1487 (°C) seem to be only related with the melting of TiSi+TiSi2 alloys.

The data on the Mo-Si phase diagram are assessed by [1991Gok], resulting in a version close to that of

[1970Sve2, 1971Sve1]. Resent publications require to insert corrections in 2 points. The first, [2000Ros]

found that the Mo3Si phase is off-stoichiometric having a homogeneity range from 23.5 to 24.0 at.% Si as

determined by EPMA and from the lattice parameters (Table 1). And secondly, [1998Fra] showed based on

their own and literature data, that only the tetragonal polymorph of MoSi2 ( -MoSi2) is stable in the binary.

Thermodynamic reassessment of [2000Liu], applied to the Mo-Si-Ti optimization in [2003Yan], shows

systematically higher temperatures for the three-phase invariant equilibria.

So, the binary systems used are: Mo-Ti [1987Mur1], Si-Ti [1996Sei] and Mo-Si after [1991Gok] corrected

for the Mo3Si and MoSi2 phases, as well as after [2000Liu].

Solid Phases

All stable solid phases are listed in Table 1. This table includes two ternary phases, (Ti1-xMox)Si2 of CrSi2crystal structure type, with 0.25 < x < 0.9 [1972Sve1, 1972Sve2, 1974Sve, 1992Boe, 1998Fra] or

0.13 < x < 0.9 [2001Wei] and a phase called T phase of unknown structure type

((32-33.75)Ti-(20.5-22.5)Mo-(45-46.1)Si (at.%) [1972Sve2]), modelled as (Ti,Mo)11Si9 by [2003Yan]

with homogeneity range from about Ti29Mo26Si45 to Ti35Mo20Si45).

Data on solubilities of the third component in the binary phases are presented in Table 2. There are

significant discrepancies between different authors, especially concerning the Ti5Si3, Ti5Si4 and Mo5Si3silicides. For these three phases the data of [2003Yan] should be preferred. The homogeneity ranges along

constant Si concentration may be strongly influenced by impurities. [1954Sch] reported a phase of Mo with

about 40 at.% Si and 2.5 at.% C to be of the Mn5Si3 type structure of Ti5Si3 and speculated on continuous

solubility between this C-stabilized modification of Mo5Si3 and Ti5Si3. The same authors claimed that

Ti5Si3 at the composition Ti3Mo2Si3 exhibits definitely selective occupation of the crystallographic

position 4(d) of space group P41212 by Mo atoms. Such an ordering was not confirmed later, furthermore

the composition Ti3Mo2Si3 was found to be outside the composition range of Ti5Si3, either at the Ti rich

end of the homogeneity range of Mo5Si3 [2003Yan] or in the two-phase field Ti5Si3+Mo5Si3 [1972Sve2].


In the system there are two pseudobinary sections Ti5Si3-Mo5Si3 and TiSi2-MoSi2.

For the former, no phase diagram is presented in literature. From the thermodynamic dataset of [2003Yan],

however, it can be calculated. The result is shown in Fig. 1. The boundaries of the two-phase field

Ti5Si3+Mo5Si3 in this figure have to be taken as tentative. Experimental data on mutual solubilities of

Ti5Si3 and Mo5Si3 are given in Table 2.

387


MSIT®

Mo–Si–Ti

[1954Sch] studied samples (Ti1-xMox)5Si3 where x = 0.15, 0.4 and 0.6, annealed at 2000° for 2 h and

reported the sample (Ti0.6Mo0.4)5Si3 to be single phase Ti5Si3, and the sample (Ti0.4Mo0.6)5Si3 to be

heterogeneous.

The TiSi2-MoSi2 pseudobinary section was studied in detail (Fig. 2). [1956Kud] presented the TiSi2-MoSi2pseudobinary system as the boundary of the pseudoternary systems TiSi2-CrSi2-MoSi2 and

TiSi2-TaSi2-MoSi2 by 1300°C isothermal sections and liquidus projections. After these authors the

Lp+(MoSi2) (Ti,Mo)Si2 (p5) reaction takes place at ~1750°C and a liquid composition of about 17 mol%

MoSi2. The equilibrium between liquid, TiSi2 and (Ti,Mo)Si2 was labelled as eutectic.

The data of Svechnikov et al. for this section [1971Koc1, 1972Sve1, 1972Sve2, 1974Sve] are mainly

consistent with the other ones, except the area between 1900 and 2000°C on the Mo-rich side, where they

assumed the hexagonal C40 phase to be stable until pure MoSi2. [1998Fra], however, showed, that the C40

structure in pure MoSi2 exists only as metastable phase after rapid solidification.

[1991Fra, 1992Boe, 1998Fra] substantiated that both phases, ternary hexagonal C40 (Ti,Mo)Si2 and

orthorhombic C54 TiSi2, are formed by peritectic reactions, Lp+MoSi2 (Ti,Mo)Si2 (p5) and

Lp+(Ti,Mo)Si2 TiSi2 (p7). The temperatures were estimated to be 1837 50 and 1513°C, respectively. The

calculation from the dataset of [2003Yan] gives close temperatures, 1861 and 1500°C, respectively. The

binary molybdenum disilicide dissolves about 1.5 mol% TiSi2. The homogeneity range of (Ti,Mo)Si2 was

measured by [1992Boe] to shrink markedly with decreasing temperature.

[1998Yi] prepared a Ti8Mo26Si66 alloy by arc melting from rod Ti of 98 mass% purity, bulk Mo of research

grade and single crystal Si (99.999999 %). The as-cast sample examined by SEM/EPMA (EDS) and TEM

was found to be two-phase, hexagonal (Ti,Mo)Si2 containing 45 6 mol% TiSi2 and orthorhombic TiSi2containing 2.1 0.6 mol% MoSi2.

[2001Wei] corroborated the previous data except the homogeneity range of (Ti,Mo)Si2, which was found

to remain significantly more extended at lower temperatures than reported by [1992Boe].

It should be noted, as the homogeneity ranges of the three phases TiSi2, (Ti,Mo)Si2 and MoSi2 are markedly

restricted to the 66.67 at.% Si line, solidus temperatures of the three-phase maxima are eventually found too

low, if the sample deviates from the exact 66.67 at.% Si line. Thus the higher solidus temperatures should

be preferred.


All invariant equilibria containing liquid were calculated and tabulated by [2003Yan]. A recalculation using

the reported thermodynamic parameters of these authors differs by 0.2 to 5°C with the tabulated

temperatures. The recalculation further shows that the reported dataset implies additional invariant

four-phase reactions in solid state. A pseudobinary eutectoid decomposition of (Ti1-xMox)Si2 into TiSi2 and

MoSi2 at 1214°C and a ternary eutectoid decomposition of (Ti1-xMox)11Si9 into Ti5Si4, Mo5Si3 and MoSi2at 1091°C may be artefacts due to not well assessed temperature dependencies of the Gibbs energies of these

two phases. For the (Ti1-xMox)Si2 phase this may be due to an overestimation of the large decrease of

solubility of TiSi2 in this phase with decreasing temperature, concluded from a single point by [1992Boe].

[2001Wei], who was not referenced by [2003Yan], reported a much smaller decrease of this solubility vs

decreasing temperature. For the (Ti1-xMox)11Si9 phase there are not sufficient experimental data to

determine well the temperature dependence of the Gibbs energy. In Table 3 the recalculated ternary

invariant (three-phase and four-phase) equilibria above 1335°C are summarized giving temperatures and

phase compositions. Equilibria calculated below 1335°C all are connected with the questionable

decompositions of the ternary phases and therefore excluded from the table. If some of the parameters are

changed by only few J mol-1, the three equilibria U7, U8 and U9 merge to only two invariant equilibria, U7*

and U8*, having together the same three-phase fields at upper and lower temperatures. These two equilibria

are tabulated by [2003Yan], they and U2 are the only ones where the temperature differs more than 2°C

between tabulated and recalculated equilibria. Figure 3 shows the reaction scheme constructed from these

recalculated invariant equilibria, except U7, U8 and U9, which are replaced by U7* and U8

*, the two

equilibria tabulated by [2003Yan].

388


MSIT®

Mo–Si–Ti


The calculated liquidus surface projection is shown in Fig. 4.

For the TiSi2-Si-MoSi2 subsystem there is an alternative version of the phase diagram experimentally

constructed by [1972Sve2] from the data on 52 ternary alloys. Instead of the calculated transition reaction

U14 at 1340°C [1972Sve2] assumed a eutectic reaction at 1320 10°C. This difference may be tolerated as

still being within the experimental accuracy. The Mo contents of liquid in the invariant equilibria given by

[1972Sve2] are significantly higher than calculated by [2003Yan]. The curvatures of the isotherms drawn

by [1972Sve2] are thermodynamically very unlikely. On the other hand data of [1974Vas] support the

eutectic reaction found by [1972Sve2]. [1974Vas] reported that the alloys with 15 to 40 mass% Ti and 5 to

20 mass% Mo (Si to balance), which were arc melted from high purity components, contain the ternary

TiSi2 + (Ti,Mo)Si2 + (Si) eutectic. Their EPMA data show that the primary disilicides TiSi2 contain 1.5-2.5

mass% Mo (i.e. 1.7 to 2.8 mol% MoSi2). The alternative liquidus surface projection from [1972Sve2] is

presented in Fig. 5.

Isothermal Sections

The 1600 and 1425°C isothermal sections were calculated by [2003Yan] (Figs. 6 and 7). The 1425°C

isothermal section drawn by [1972Sve2] disagrees mainly by very small extensions of the ternary

homogeneity ranges except Ti5Si3 and Mo5Si3. As a consequence [2003Yan] calculated equilibria between

Ti5Si4 and Mo5Si3, Ti5Si4 and TiSi2, Ti5Si4 and (Ti,Mo)Si2, whereas [1972Sve2] reported equilibria

between Ti5Si3 and (Ti,Mo)11Si9, TiSi and (Ti,Mo)11Si9. In spite of the large extensions of the homogeneity

ranges of Ti5Si3 and Mo5Si3 [1972Sve2] placed the corners of the three-phase fields of these phases with

(Ti,Mo)11Si9 and the next phase richer in Si very near to the binary boundaries which is thermodynamically

very unlikely. The TiSi2-Si-MoSi2 isothermal section at 1250°C was experimentally studied in [1972Sve]

(Fig. 8).

The early results of [1956Bre] as well as of [1965Gar] are well consistent with those of [1972Sve2] and

[2003Yan] as far as equilibrium was reached.

[2002Str] studied an alloy Mo3Ti2Si3 (18.6Ti-62.4Mo-18.7Si-0.022C-0.014N-0.012O (mass%), from

chemical analysis). The microstructural and phase analyses were carried out using SEM/EMPA (EDS) and

powder XRD experiments. The alloy annealed at 1600°C for 120 h in Ar was turned out to be two-phase,

Mo5Si3 as the major phase and Ti5Si3. That agrees well with the isothermal section of [2003Yan] (Fig. 5).


[1972Sve2] published 8 vertical sections for the TiSi2-Si-MoSi2 subsystem at constant mass contents of Si

of 80, 75, 70, 65, 60, 55, 50 and 45 mass% (Fig. 9), which are revised for the existence of the CrSi2 crystal

type phase only as ternary (Ti1-xMox)Si2 (0.25 < x < 0.9). The 65 mass% Si section was assumed in

[1972Sve2] to pass through the invariant points E and U13, therefore the two-phase field L+(Si) or

L+(Ti,Mo)Si2 is narrowed almost to a line (Fig. 9d). However, these vertical sections disagree significantly

with the corresponding ones calculated from the dataset of [2003Yan]. Whether this is due to the dataset or

due to the conclusions from the experimental data is not clear; here reported (Fig. 9) are the experimental

diagrams.

Thermodynamics

Thermodynamic parameters optimized by [2003Yan] are presented in Table 4 together with the adopted

binary parameters of [1996Sei, 1998Sau, 2000Liu].


[1992Boe] found the fracture toughness and oxidation resistance (at 1300°C for 10 h in air) of MoSi2alloyed with 3 to 18 mol% TiSi2 to be similar to those of unalloyed MoSi2. The samples were two-phase

(Ti,Mo)Si2+MoSi2 of equiaxed morphology.

389


MSIT®

Mo–Si–Ti

[2002Nak] prepared lamellar structure in (Ti,Mo)Si2+(MoSi2) alloys by zone melting. The alloys

(Ti1-xMox)Si2 (where x = 0.50 to 0.95) were arc-melted from Ti of 99.99 mass% purity, Mo of 99.9 mass%

and Si of 99.9999 mass% and studied by XRD and TEM. After zone melting the samples were annealed at

1400°C for 24 h. In the whole rod the structure was lamellar for the (Ti0.30Mo0.70)Si2 alloy.

Oxidation experiments were carried out by [1995Yan] in air (both isothermal between 1434 and 1685°C

and cyclic from room temperature to 1630°C tests) for the alloy (Ti0.1Mo0.9)Si2. The disilicide was found

to form a good protective scale up to 1550°C (the Cristobalite SiO2 “plus” Rutile TiO2 eutectic) at

isothermal conditions, which however had many cracks and case spalling after the cyclic oxidation.

Miscellaneous

The pesting behavior of (Ti0.1Mo0.9)Si2 alloys was examined by [1996Yan] at 400 to 700°C in air.

Silicide formation and resistivity were studied by [1997Cab] for Ti alloys containing to 20 at.% Mo

deposited on polycrystalline Si.

[2003Her] studied a mechanical alloying process experimentally and developed a computer simulation for

the powder mixture of 16.7Ti - ~27Mo-Si (mass%) (to balance). The hexagonal disilicide content was found

to be accumulated and the amount of tetragonal form reaches maximum and goes down.

References

[1952Now] Nowotny, H., Kieffer, R., Schachner, H., “Structure Investigations in Disilisides” (in

German), Monatsh. Chem., 83, 1243-1252 (1952) (Equi. Diagram, Crys. Structure,

Experimental, 8)

[1954Sch] Schachner, H., Cerwenka, E., Nowotny, H., “New Silicides of M5Si3-Type with D88

Structure” (in German), Monatsh. Chem., 85, 245-254 (1954) (Equi. Diagram, Crys.


[1956Bre] Brewer, L., Krikorian, O., “Reactions of Refractory Silicides with Carbon and Nitrogen”,

J. Electrochem. Soc., 103(1), 38-51 (1956) (Crys. Structure, Experimental, 52)

[1956Kud] Kudielka, H., Nowotny, H., “Disilicide Systems” (in German), Monatsh. Chem., 87(3),

471-482 (1954) (Equi. Diagram, Crys. Structure, Experimental, 8)

[1965Gar] Garfinkle, M., Davis, H.M., “Reaction of Liquid Ti with Some Refractory Compounds”,

Trans. Qart. ASM, 58, 520-530 (1965) (Crys. Experimental, Equi. Diagram, 9)

[1970Sve1] Svechnikov, V.N., Kocherzhynsky, Yu.A., Yupko, L.M., “Phase Diagram

Titanium-Silicon” (in Russian), Doklady Akad. Nauk SSSR, 193(2), 393-396 (1970) (Equi.

Diagram, Crys. Structure, Experimental, 16)

[1970Sve2] Svechnikov, V.M., Kocherzhinsky, Yu.O., Yupko, L.M., “Phase Diagram

Molybdenum-Silicon” (in Ukrainian), Dopov. Akad. Nauk Ukrain. RSR, A(6), 553-557

(1970) (Equi. Diagram, Crys. Structure, Experimental, 14)

[1971Koc1] Kocherzhinsky, Yu.A., “Differential Thermocouple up to 2450°C and Thermographic

Investigation of Refractory Silisides”, Therm. Anal., Proc. Third ICTA (Davos), 1, 549-559


[1971Koc2] Kocherzhinsky, Yu.A., Shishkin, Ye.A., Vasilenko, V.I., “Apparatus for Differential

Thermal Analysis with Temperature-Sensing Thermocouple up to 2200°C” (in Russian),

Ageev, N.V., Ivanov, O.S. (Eds.), Nauka, Moskow, 245-249 (1971) (Equi. Diagram, Crys.


[1971Sve1] Svechnikov, V.M., Kocherzhinsky, Yu.O., Yupko, L.M., “Phase Diagram

Molybdenum-Silicon”, in “Diagrammy Sostoyaniya Metal. System” (in Russian), Ageev,

N.V., Ivanov, O.S. (Eds.), Nauka, Moskow, 116-119 (1971) (Equi. Diagram, Crys.


[1971Sve2] Svechnikov, V.M., Kocherzhinsky, Yu.O., Yupko, L.M., Kulik, O.G., Shishkin, Ye.A.,

“Phase Diagram Titanium-Silicon”, in “Diagrammy Sostoyaniya Metal. System” (in

Russian), Ageev, N.V., Ivanov, O.S., (Eds.), Nauka, Moskow, 129-135 (1971) (Equi.

Diagram, Crys. Structure, Experimental, 22)

390


MSIT®

Mo–Si–Ti

[1972Sve1] Svechnikov, V.N., Kocherzhynsky, Yu.A., Yupko, L.M., “Phase Diagram MoSi2-TiSi2” (in

Russian), Dokl. Akad. Nauk SSSR, 206(3), 631-633 (1972) (Equi. Diagram, Crys. Structure,

Experimental, 10)

[1972Sve2] Svechnikov, V.M., Kocherzhinsky, Yu.O., Yupko, L.M., “Investigation of Phase Equilibria

in Alloys of Silicon with Molybdenum and Titanium” (in Ukrainian), Dop. Akad. Nauk

Ukrain. RSR, A(6), 566-570 (1972) (Equi. Diagram, Crys. Structure, Experimental, 5)

[1974Sve] Svechnikov, V.N., Kocherzhinsky, Yu.A., Yupko, L.M., “Phase Diagrams of the Systems

MoSi2-TiSi2 and MoSi2-CrSi2”, in “Structure of Phases, Phase Transformation and Phase

Diagrams for Metal Systems” (in Russian), Ivanov, O.S., Alekseeva, Z.M., (Eds.), Nauka,

Moskow, 132-136 (1974) (Equi. Diagram, Crys. Structure, Experimental, 14)

[1974Vas] Vasilyeva, E.V., Terentyeva, V.S., ”Study of the Structure, Phase Composition, and Heat

Resistance of Alloys” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 10(6), 1023-1026


[1987Mur1] Murray, J.L., “The Mo-Ti (Molybdenum-Titanium) System”, in “Phase Diagrams of Binary

Titanium Alloys”, Murray, J.L. (Ed.), ASM, Metals Park, Ohio, 169-175 (1987) (Equi.

Diagram, Crys. Structure, Thermodyn., Assessment, 45)

[1987Mur2] Murray, J.L., “The Si-Ti (Silicon-Titanium) System”, in “Phase Diagrams of Binary

Titanium Alloys”, Murray, J.L. (Ed.), ASM, Metals Park, Ohio, 289-293 (1987) (Equi.

Diagram, Crys. Structure, Thermodyn., Assessment, 29)

[1991Din] Dinsdale, A.T., “SGTE Data for Pure Elements”, Calphad, 15(4), 317-425 (1991)

(Thermodyn.)

[1991Fra] Frankwisc, P.S., Perepezko, J.H., “Phase Stability and Solidification Pathways in MoSi2Based Alloys”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV,

213, 169-174 (1991) (Equi. Diagram, Crys. Structure, Experimental, 6)

[1991Gok] Gokhale, A.B., Abbaschian, G.J., “The Mo-Si (Molibdenum-Silicon) System”, J. Phase

Equilib., 12(4), 493-498 (1991) (Equi. Diagram, Crys. Structure, Thermodyn.,

Assessment, 36)

[1992Boe] Boettinger, W.J., Peperezko, J.H., Frankwicz, P.S., “Application of Ternary Phase

Diagrams to the Development of MoSi2-Based Materials”, Mater. Sci. Eng. A, A155(1),

33-44 (1992) (Equi. Diagram, Crys. Structure, Phys. Prop., Experimental, 18)

[1995Mae] Maex, K., Van Rossum, M., Reader, A., “Crystal Structure of TM Silicides”, Emis Datarev.

Ser. 14 (Properties of Metal Silicides), 3-14 (1995) (Crys. Structure, Review, 102)

[1995Yan] Yanagihara, K., Maruyama, T., Nagata, K., “High Temperature Oxidation of Mo-Si-X

Intermetallics (X = Al, Ti, Ta, Zr and Y)”, Intermetallics, 3(3), 243-251 (1995) (Phys. Prop.,

Experimental, 22)


Z. Metallkd., 87, 2-13 (1996) (Equi. Diagram, Thermodyn., Assessment, 63)

[1996Yan] Yanagihara, K., Maruyama, T., Nagata, K., “Effect of Third Elements on the Pesting

Suppression of Mo-Si-X Intermetallics (X = Al, Ta, Ti, Zr and Y)”, Intermetallics, 4(1),

S133-S139 (1996) (Crys. Structure, Phys. Prop., Experimental, 31)

[1997Cab] Cabral, C.Jr., Clevenger, L.A., Harper, J.M.E., d’Heurle, F.M., Roy, R.A., Lavoie, C.,

Saenger, K.L., Miles, G.L., Mann, R.W., Nakos, J.S., “Low Temperature Formation of

C54-TiSi2 Using Titanium Alloys”, Appl. Phys. Lett., 71(24), 3531-3533 (1997)

(Experimental, 13)

[1998Du] Du, Y., Schuster, J.C., “A Reinvestigation of the Constitution of the Partial System TiSi-Si”,

J. Mater. Sci. Lett., 17, 1407-1408 (1998) (Equi. Diagram, 7)

[1998Fra] Frankwicz, P.S., Perepezko, J.H., “Phase Stability of MoSi2 in the C11b and C40 Structures

at High Temperatures”, Mater. Sci. Eng. A, A246(1-2), 199-206 (1998) (Equi. Diagram,

Experimental, 55)

[1998Sau] Saunders, N., “System Mo-Ti” in “COST 507 Thermochemical Database for Light Metal

Alloy”, Ansara, A.T. Dinsdale, M.H. Rand (Eds.), European Communities, Luxembourg,

Vol. 2, 249-252 (1998) (Thermodyn, Assessment, 0)

391


MSIT®

Mo–Si–Ti

[1998Yi] Yi, D., Lai, Z., Li, Ch., Akselsen O.M., Ulvensoen, J.H., “Ternary Alloying Study of

MoSi2”, Metall. Mater. Trans. A, A29(1), 119-129 (1998) (Experimental, 29)

[2000Liu] Liu, Y., Shao, G., Tsakiropoulos, P., “Thermodynamic Reassessment of the Mo-Si and

Al-Mo-Si Systems”, Intermetallics, 8(8), 953-962 (2000) (Equi. Diagram, Thermodyn.,

Assessment, 48)

[2000Ros] Rosales, I., Schneibel, J.H., “Stoichiometry and Mechanical Properties of Mo3Si”,

Intermetallics, 8, 885-889 (2000) (Crys. Structure, Experimental, 20)

[2000Wil] Williams, J.J., Kramer, M.J., Akinc, M., Malik, S.K., “Effects of Interstitial Additions on

the Structure of Ti5Si3”, J. Mater. Res., 15(8), 1773-1779 (2000) (Crys. Structure,

Experimental, 16)

[2001Tan] Tanaka, K., Nawata, K., Inui, H., Yamaguchi, M., Koiwa, M., “Refinement of

Crystallographic Parameters in Tansition Metal Disilicides with the C11b, C40 and C54

Structures”, Intermetallics, 9, 603-607 (2001) (Crys. Structure, Experimental, 12)

[2001Wei] Wei, F.-G., Kimura, Y., Mishima, Y., “Microstructure and Phase Stability in MoSi2-TSi2(T = Cr, V, Nb, Ta, Ti) Pseudo-Binary Systems”, Mater. Trans., JIM, 42(7), 1349-1355


[2002Nak] Nakano, T., Nakai, Y., Maeda, S., Umakoshi, Y., “Microstructure of Duplex-Phase

NbSi2(C40)/MoSi2(C11b) Crystals Containing a Single Set of Lamellae”, Acta Mater.,

50(7), 1781-1795 (2002) (Crys. Structure, Experimental, 38)

[2002Str] Ström, E., Zhang, J., Eriksson, S., Li, CH., Feng, D., “The Influence of Alloying Elements

on Phase Constitution and Microstructure of Mo3M2Si3 (M = Cr, Ti, Nb, Ni or Co)”, Mater.

Sci. Eng. A, A329-331, 289-294 (2002) (Equi. Diagram, Experimental, 12)

[2003Her] Heron, A.J., Schaffer, G.B., “Mechanical Alloying of MoSi2 with Ternary Alloying

Elements. Part 2. Computer Simulation”, Mater. Sci. Eng. A, A359(1-2), 319-325 (2003)


[2003Yan] Yang, Y., Chang, Y.A., Tan, L., Du, Y., “Experimental Investigation and Thermodynamic

Descriptions of the Mo-Si-Ti System”, Mater. Sci. Eng. A, A361(1-2), 281-293 (2003)

(Equi. Diagram, Thermodyn., Experimental, Assessment, 15)

[2004Yan] Yang, Y., Chang, Y.A., Private communication (2004)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.03

c = 468.10

[Mas2, 1996Sei]

dissolves up to 0.54 at.% Si

, ( Ti,Mo)

< 2623

cI2

Im3m

W

a = 331.12

a = 314.51

pure Ti [Mas2]

pure Mo [Mas2]

Ti and Mo dissolves up to 5 and 4 at.%

Si, respectively [1991Gok, 1996Sei]

(Si)

< 1414

cF8

Fm3m

C (diamond)

a = 543.09 [Mas2]

Ti3Si

< 1170

tP32

P42/n

Ti3P

a = 1019.6

c = 509.7

[1996Sei, V-C2]

392


MSIT®

Mo–Si–Ti

* Taken from a figure; parameter c is not reported.

Ti5Si3< 2130

hP16

P63/mcm

Mn5Si3

a = 746.0 0.2

c = 515.2 0.2

a = 733*

35.7-40 at.% Si

[1996Sei, 2000Wil]

for (Ti0.65Mo0.35)5Si3 [1972Sve2]

Ti5Si4< 1920

tP36

P41212

Zr5Si4

a = 713.3

c = 1297.7

[1996Sei, 1970Sve1, 1971Sve2]

TiSi

< 1570

oP8 or oP8

Pmm2 Pnma

TiSi FeB

a = 654.4

b = 363.8

c = 499.75

[1996Sei, 1995Mae]

TiSi2< 1500

oF24

Fddd

TiSi2

a = 826.80 0.03

b = 480.02 0.01

c = 855.21 0.06

[1987Mur2, 2001Tan]

Mo3Si

< 2025

cP8

Pm3n

Cr3Si

a = 489.6

a = 489.0

23.5-24 at.% Si

[1991Gok, 2000Ros]

Mo5Si3< 2180

tI32

I4/mcm

W5Si3

a = 964.83 0.06

c = 491.35 0.05

a = 975*

37-40 at.% Si

[1991Gok, 1995Mae]

for (Ti0.4Mo0.6)5Si3 [1972Sve2]

MoSi2< 2020

tI6

I4/mmm

MoSi2

a = 320.56 0.03

c = 784.50 0.04

[1991Gok, 2001Tan]

-MoSi2metastable in the

binary

(Ti1-xMox)Si2(0.25 < x < 0.9)

< 850

hP9

P6222 or P6422

CrSi2

a = 464.2 0.5

c = 652.9 0.5

464.5*< a < 470.6

a = 467.4 0.5

c = 650.2 0.5

[1998Fra, 1970Sve2, 1971Sve1]

at 0.3 < x < 0.9

at Mo0.6Ti0.4Si2[1972Sve1, 1974Sve]

T phase

1935

(Ti,Mo)11Si9

? ? 45-46.1 at.% Si [1972Sve2]

20.5-22.5 at.% Mo [2003Yan]

at 45 at.% Si

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

393


MSIT®

Mo–Si–Ti

Table 2: Experimentally Determined Ternary Solubilities of the Mo-Si-Ti Phases


Phase Temperature [°C] Solubility Comments/Reference

, ( Ti,Mo) 1425

1600

Ti86.4Mo9.0Si4.6

Ti68.0Mo28.5Si3.5

[2003Yan]

[2003Yan]

Ti5Si3 1425

1425

1600

1800

30 mol% Mo5Si319 mol% Mo5Si320.5 mol% Mo5Si335 mol% Mo5Si3

[1972Sve2]

[2003Yan]

[2003Yan]

[1972Sve2]

Ti5Si4 1425

1425

1600

1.7 mol% “Mo5Si4”

24 mol% “Mo5Si4”

19 mol% “Mo5Si4”

[1972Sve2]

[2003Yan]

[2003Yan]

TiSi 1425 2 mol% “MoSi” [1972Sve2]

TiSi2 1500

1200

1510 10

1400

1400

13 mol% MoSi218 mol% MoSi2~10 mol% MoSi2~5 mol% MoSi2~5 mol% MoSi2

[1972Sve2]

[1972Sve2]

[1992Boe]

[1992Boe]

[2001Wei]

Mo3Si 1425

1600

39 mol% Ti3Si

41 mol% Ti3Si

[2003Yan]

[2003Yan]

Mo5Si3 1425

1425

1600

1800

41 mol% Ti5Si366.7 mol% Ti5Si365.5 mol% Ti5Si341 mol% Ti5Si3

[1972Sve2]

[2003Yan]

[2003Yan]

[1972Sve2]

MoSi2 1850 ~1.5 mol% TiSi2 [1992Boe]

(Ti1-xMox)Si2 1300

1700

1500

1300

1850

1600

1500

1400

1400

0.07 x 0.6

0.11 x

0.125 x 0.74

0.14 x 0.68

0.14 x

0.23 x

x 0.75

0.28 x 0.59

0.21 x 0.85

[1952Now]

[1972Sve2]

[1972Sve2]

[1972Sve2]

[1992Boe]

[1992Boe]

[1992Boe]

[1992Boe]

[2001Wei]


Mo Si Ti

L Mo5Si3 + Ti5Si3 2207.9 e1 L

Mo5Si3Ti5Si3

18.1

24.8

17.6

37.5

37.4

37.5

44.4

37.8

44.9

L +Mo5Si3 Mo3Si 2126.3 p1 L

Mo5Si3Mo3Si

58.9

41.8

58.8

24.9

36.9

16.2

16.2

11.3

25.0

394


MSIT®

Mo–Si–Ti

L ( Ti,Mo) + Mo3Si 2123.0 e2 L

( Ti,Mo)

Mo3Si

62.9

93.4

59.7

22.9

2.5

15.3

14.2

4.1

25.0

L +Mo5Si3 (Ti,Mo)11Si9 1957.1 p3 L

Mo5Si3(Ti,Mo)11Si

10.7

28.4

37.6

52.7

37.6

17.4

36.6

34.0

45.0

L +Mo5Si3 Ti5Si3 +

(Ti,Mo)11Si9

1937.0 U1 L

Mo5Si3Ti5Si3(Ti,Mo)11Si9

5.8

32.1

14.9

16.5

51.4

37.8

37.5

45.0

42.8

39.3

47.6

38.5

L +Mo5Si3 Ti5Si3 + Mo3Si 1935.9 U2 L

Mo5Si3Ti5Si3Mo3Si

29.9

23.3

42.5

13.9

19.7

36.2

25.0

37.5

50.4

40.5

32.5

48.6

L +Mo3Si ( Ti,Mo) + Ti5Si3 1930.4 U3 L

Mo3Si

( Ti,Mo)

Ti5Si3

29.6

42.4

68.9

13.7

19.6

25.0

1.8

37.5

50.8

32.6

29.3

48.8

L + Ti5Si3 Ti5Si4 +

(Ti,Mo)11Si9

1916.3 U4 L

Ti5Si3Ti5Si4(Ti,Mo)11Si9

4.8

13.7

9.3

15.7

51.6

37.5

44.4

45.0

43.6

48.8

46.3

39.3

Ti5Si3 + (Ti,Mo)11Si9 Mo5Si3+ Ti5Si4

1870.8 U5 Ti5Si3(Ti,Mo)11Si9Mo5Si3Ti5Si4

14.5

16.2

22.8

9.7

37.5

45.0

37.6

44.4

48.0

38.8

39.6

45.9

Ti5Si3 + Mo3Si ( Ti,Mo) +

Mo5Si3

1870.2 U6 Ti5Si3Mo3Si

( Ti,Mo)

Mo5Si3

13.5

42.5

69.7

23.0

37.5

25.0

1.5

36.2

49.0

32.5

28.8

40.8

L + MoSi2 (Ti,Mo)Si2 1860.8 p5 L

(Ti,Mo)Si2

19.4

13.9

66.7

66.7

13.9

19.4

L + MoSi2 Mo5Si3 +

(Ti,Mo)Si2

1775.1 U7 L

Mo5Si3(Ti,Mo)Si2

21.4

41.0

29.1

59.6

39.1

66.7

19.0

19.9

4.2

L + Mo5Si3 (Ti,Mo)Si2 +

(Ti,Mo)11Si9

1772.9 U8 L

Mo5Si3(Ti,Mo)Si2(Ti,Mo)11Si9

21.2

40.8

29.0

28.0

59.6

39.1

66.7

45.0

19.2

20.1

4.3

27.0

Mo5Si3 + (Ti,Mo)Si2(Ti,Mo)11Si9 + MoSi2

1768.4 U9 Mo5Si3(Ti,Mo)Si2(Ti,Mo)11Si9

40.7

29.0

28.0

39.1

66.7

45.0

20.2

4.3

27.0

L + (Ti,Mo)11Si9 Ti5Si4 +

(Ti,Mo)Si2

1555.3 U10 L

(Ti,Mo)11Si9Ti5Si4(Ti,Mo)Si2

3.4

17.8

10.7

15.8

62.9

45.0

44.4

66.7

33.7

37.2

44.8

17.5


Mo Si Ti

395


MSIT®

Mo–Si–Ti

Table 4: Binary and Ternary Thermodynamic Parameters* for the Mo-Si-Ti Phases.

L + (Ti,Mo)Si2 TiSi2 1499.9 p7 L

(Ti,Mo)Si2TiSi2

1.2

9.5

2.6

66.7

66.7

66.7

32.1

23.8

30.7

L + (Ti,Mo)Si2 TiSi2 + Ti5Si4 1491.6 U11 L

(Ti,Mo)Si2TiSi2Ti5Si4

1.4

9.9

2.6

6.1

63.7

66.7

66.7

44.4

34.9

23.4

30.7

49.5

L + Ti5Si4 TiSi + TiSi2 1491.2 U12 L

Ti5Si4TiSi2

1.3

5.9

2.6

63.7

44.4

66.7

35.0

49.7

30.7

L + MoSi2 (Ti,Mo)Si2 + (Si) 1375.3 U13 L

(Ti,Mo)Si2

2.5

23.7

87.5

66.7

10.0

9.6

L + (Ti,Mo)Si2 TiSi2 + (Si) 1340.4 U14 L

(Ti,Mo)Si2

1.3

16.7

83.0

66.7

15.7

16.6

Phase Parameter References

Liquid 0LMo,Siliq

1LMo,Siliq

2LMo,Siliq

0LMo,Tiliq

0LTi,Siliq

1LTi,Siliq

2LTi,Siliq

LMo,Si,Tiliq

-146600.+16.5 T

7100.-3. T

8000.+16.5 T

-9000.+2. T

-242611.79+17.57906 T

26753.36-2.14028 T

76881.70-6.15055 T

71636.04 xMo-292917.61

xSi+229440.02 xTi

[2000Liu]

[2000Liu]

[2000Liu]

[1998Sau]

[1996Sei]

[1996Sei]

[1996Sei]

[2000Liu]

( Ti) 0LMo,Ti

0LSi,Ti

1LSi,Ti

2LSi,Ti

22760.-6. T

-302072.20+77.76859 T

26753.36-2.14028 T

76881.70-6.15055 T

[1998Sau]

[1996Sei]

[1996Sei]

[1996Sei]

, ( Ti,Mo) 0LMo,Si

0LMo,Ti

1LMo,Ti

0LSi,Ti

1LSi,Ti

2LSi,Ti

-60000.+5. T

2000.

-2000.

-262502.74+38.81506 T

26753.36-2.14028 T

76881.70-6.15055 T

[2000Liu]

[1998Sau]

[1998Sau]

[1996Sei]

[1996Sei]

[1996Sei]

Ti3Si 0GTi:SiTi3Si - 30GTi

hcp - 0GSidiamond -200000.+4.38996 T [1996Sei]


Mo Si Ti

396


MSIT®

Mo–Si–Ti

* The unary parameters 0GMobcc, 0GMo

liq, 0GSidiamond, 0GSi

liq, 0GTihcp, 0GTi

bcc and 0GTiliq have to be taken from

[1991Din]

Ti5Si30GMo:Si

Ti5Si3 - 50GMobcc - 30GSi

diamond

0GTi:SiTi5Si3 - 50GTi

hcp - 30GSidiamond

0LMo,Ti:SiTi5Si3

-280000.

-579282.48+2.32 T

-295990.88+64.608 T

[2003Yan]

[1996Sei]

[2003Yan]

Ti5Si40GMo:Si

Ti5Si4 - 50GMobcc - 40GSi

diamond

0GTi:SiTi5Si4 - 50GTi

hcp - 40GSidiamond

0LMo,Ti:SiTi5Si4

-350000.

-711000.+26.79 T

-94929.06

[2003Yan]

[1996Sei]

[2003Yan]

TiSi 0GTi:Si TiSi - 0GTi

hcp - 0GSidiamond -156888.96+9.74346 T [1996Sei]

TiSi20GMo:Si

TiSi2 - 0GMobcc - 20GSi

diamond

0GTi:SiTiSi2 - 0GTi

hcp - 20GSidiamond

0LMo,Ti:SiTiSi2

-69872.53

-175619.85+7.401 T

-166771.45-26.54 T

[2003Yan]

[1996Sei]

[2003Yan]

Mo3Si 0GMo:SiMo3Si - 30GMo

bcc - 0GSidiamond

0GTi:SiMo3Si - 30GTi

hcp - 0GSidiamond

0LMo,Ti:SiMo3Si

1LMo,Ti:SiMo3Si

-111600.-1.12 T

-144000.

-213756.36+27.12 T

-124904.32+47.88 T

[2000Liu]

[1996Sei]

[2003Yan]

[2003Yan]

Mo5Si30GMo:Si

Mo5Si3 - 50GMobcc - 30GSi

diamond

0GSi:SiMo5Si3 - 80GSi

diamond

0GTi:SiMo5Si3 - 50GTi

hcp - 30GSidiamond

0GMo:MoMo5Si3 - 80GMo

bcc

0GSi:MoMo5Si3 - 30GMo

bcc - 50GSidiamond

0LMo,Si:MoMo5Si3

0LMo,Si:MoMo5Si3

0LMo,Ti:SiMo5Si3

0LMo:Mo,SiMo5Si3

0L Si:Mo,SiMo5Si3

-311840.-10.792 T

80000.

-540000.

80000.

471840.+10.792 T

80000.

-548000.+172. T

-478743.44+132.96 T

160000.

160000.

[2000Liu]

[2000Liu]

[2003Yan]

[2000Liu]

[2000Liu]

[2000Liu]

[2000Liu]

[2003Yan]

[2000Liu]

[2000Liu]

MoSi20GMo:Si

MoSi2 - 0GMobcc - 20GSi

diamond -135468.+0.669 T [2000Liu]

(Ti1-xMox)Si20GMo:Si

(Ti,Mo)Si2 - 0GMobcc - 20GSi

diamond

0GTi:Si(Ti,Mo)Si2 - 0GTi

hcp - 20GSidiamond

0LMo,Ti:Si(Ti,Mo)Si2

-121186.90-5.167 T

-125974.57-19.25956 T

0

[2003Yan]

[2003Yan]

[2003Yan]

(Ti1-xMox)11Si90GMo:Si

(Ti,Mo)11Si9 - 110GMobcc - 90GSi

diamond

0GTi:Si(Ti,Mo)11Si9 - 110GTi

hcp - 90GSidiamond

0LMo,Ti:Si(Ti,Mo)11Si9

-658372.

-1362522.

-743100.8

[2003Yan]

[2003Yan]

[2003Yan]

Phase Parameter References

397


MSIT®

Mo–Si–Ti

10 20 30 40 50 60

1800

1900

2000

2100

2200

Ti 62.50

Mo 0.00

Si 37.50

Ti 0.00

Mo 62.50

Si 37.50Mo, at.%

Te

mp

era

ture

, °C

Mo5Si3

L+Mo5Si3

Ti5Si3+Mo5Si3

L+Ti5Si3

Ti5Si3

LFig. 1: Mo-Si-Ti.

Pseudobinary system

Ti5Si3 - Mo5Si;

calculated from the

dataset of [2003Yan]

(Table 4)

30 20 10

1250

1500

1750

2000

Ti 33.30

Mo 0.00

Si 66.70

Ti 0.00

Mo 33.30

Si 66.70Ti, at.%

Te

mp

era

ture

, °C

L+(Ti,Mo)Si2

TiSi2

LL+MoSi2

(Ti,Mo)Si2

MoSi2

2020°C

1850

1510+/-10p7

Fig. 2: Mo-Si-Ti.

Pseudobinary system

TiSi2 - MoSi2; dashed

lines: calculation

[2003Yan], solid

lines: assessed from

experimental data

[1972Sve, 199Boe,

2001Wei]

398


MSIT®

Mo–Si–Ti

Fig

. 3:

M

o-S

i-T

i. R

eact

ion s

chem

e, c

alcu

late

d f

rom

the

dat

aset

of

[2003Y

an]

(Tab

le 4

); U

7* a

nd U

8*,

taken

fro

m T

able

4 o

f [2

003Y

an]

repla

ce U

7 +

U8 +

U9

of

Tab

le 3

of

this

rep

ort

Si-

Ti

Mo

-Si

Mo

-Si-

Ti

Lβ

+ M

o3S

i

21

23

e 2(m

ax)

l +

Τi 5

Si 3

Ti 5

Si 4

19

20

p4

L+

Mo

5S

i 3T

i 5S

i 3+

(Ti,

Mo

) 11S

i 91

937

U1

l +

β M

o3S

i

20

64

(20

25)

p2

l +

Τi 5

Si 4

TiS

i

15

70

p6

l T

iSi

+ T

iSi 2

14

74

e 5

lβ

+ T

i 5S

i 3

13

40

e 7

l T

iSi 2

+ (

Si)

13

30

e 8

l M

o3S

i +

Mo

5S

i 3

20

63

(20

20)

e 3

l M

o5S

i 3 +

MoS

i 2

19

23

(19

00)

e 4

l M

oS

i 2 +

(S

i)

14

07

(14

00)

e 6

L +

Mo

5S

i 3 M

o3S

i

2126.3

p1(m

ax)

L M

o5S

i 3 +

Ti 5

Si 3

2207.9

e 1(m

ax)

L+

Μo

5S

i 3(T

i,M

o) 1

1S

i 9

1957.1

p3(m

ax)

L+

Mo

5S

i 3T

i 5S

i 3+

Mo

3S

i1935.9

U2

L+

Mo

3S

iβ

+ T

i 5S

i 31930.4

U3

L+

Ti 5

Si 3

(Ti,

Mo

) 11S

i 9+

Ti 5

Si 4

1916.3

U4

L+

Mo

5S

i 3M

oS

i 2+

(Ti,

Mo

) 11S

i 91775.1

U7

*

L+

(Ti,

Mo) 1

1S

i 9(T

i,M

o)S

i 2+

Ti 5

Si 4

1555.3

U1

0

L+

(Ti,

Mo)S

i 2T

iSi 2

+T

i 5S

i 41491.6

U1

1

L +

Ti 5

Si 4

TiS

i 2+

TiS

i1491.2

U1

2

L +

MoS

i 2(T

i,M

o)S

i 2+

(S

i)1375.3

U1

3

L +

(T

i,M

o)S

i 2T

iSi 2

+ (

Si)

1340.4

U1

4

L+

Mo

Si 2

(Ti,

Mo

)Si 2

1860.8

p5(m

ax)

L+

(Ti,

Mo

)Si 2

TiS

i 2

1499.9

p7(m

ax)

L+

Mo

Si 2

(Ti,

Mo

)Si 2

+(T

i,M

o) 1

1S

i 91772.9

U8

*M

o5S

i 3+

Mo

Si 2

+(T

i,M

o) 1

1S

i 9

Mo

Si 2

+(T

i,M

o) 1

1S

i 9+

(Ti,

Mo

)Si 2

(Ti,

Mo

) 11S

i 9+

(Ti,

Mo

)Si 2

+T

i 5S

i 4

Mo

Si 2

+(T

i,M

o)S

i 2+

(Si)

(Ti,

Mo

)Si 2

+T

iSi 2

+T

i 5S

i 4

Ti 5

Si 4

+T

iSi 2

+T

iSi

(Ti,

Mo

)Si 2

+T

iSi 2

+(S

i)

Ti 5

Si 3

+(T

i,M

o) 1

1S

i 9M

o5S

i 3+

Ti 5

Si 4

1870.8

U5

Ti 5

Si 3

+M

o3S

iβ+

Mo

5S

i 31870.2

U6

Ti 5

Si 3

+M

o5S

i 3+

Ti 5

Si 4

(Ti,

Mo

) 11S

i 9+

Mo

5S

i 3+

Ti 5

Si 4

Mo

5S

i 3+

β+M

o3S

iM

o5S

i 3+

β+T

i 5S

i 3

399


MSIT®

Mo–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

MoSi2

Mo5Si

3

β

Mo3Si

Ti5Si

3

(Ti,Mo)11

Si9

(Ti,Mo)Si2

(Si)

TiSi2

TiSi

Ti5Si

4

e8

e5

p6

p4

e7

e6

e4

e3

p2

U3

U2

U4

U1

p3

U7

U8

p5

U13

U14

U11

U12

U10

e1

p1

e2

p7

1500 1600

1800

1900

2000

2100

2200

2300

2400

2500

2200

2100

20001900

18001700 1900 2000

1600

1500

2100

Fig. 4: Mo-Si-Ti.

Liquidus surface

projection, calculated

from the dataset of

[2003Yan] (Table 4)

10

20

30

10 20 30

70

80

90

Ti 33.33

Mo 0.00

Si 66.67

Ti 0.00

Mo 33.33

Si 66.67


Axes: at.%

MoSi2

TiSi2

(Si)

p5

e8

p7

e6

U13

E

1400

1400

1500

1600 17

00

1800

1900

2000

1500

(Ti,Mo)Si2

Fig. 5: Mo-Si-Ti.

Liquidus surface

projection in the

TiSi2-MoSi2-Si range

after [1972Sve2]

400


MSIT®

Mo–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

L

β

MoSi2

(Ti,Mo)Si2

Ti5Si

4

Ti5Si

3

L

Mo3Si

Mo5Si

3

(Ti,Mo)11

Si9

Fig. 6: Mo-Si-Ti.


1600°C [2003Yan]

20

40

60

80

20 40 60 80

20

40

60

80

Ti Mo


Axes: at.%

β

L

(Ti,Mo)Si2

TiSi2

MoSi2

Mo5Si

3

Mo3Si

TiSi

Ti5Si

4

Ti5Si

3

L

Fig. 7: Mo-Si-Ti.


1425°C

[2003Yan, 2004Yan]

401


MSIT®

Mo–Si–Ti

10

20

30

10 20 30

70

80

90

Ti 35.00

Mo 0.00

Si 65.00

Ti 0.00

Mo 35.00

Si 65.00


Axes: at.%

MoSi2

TiSi2

(Ti,Mo)Si2

(Si)Fig. 8: Mo-Si-Ti.

Partial isothermal

section at 1250°C

10

1200

1300

1400

1500

1600

1700

Ti 12.80

Mo 0.00

Si 87.20

Ti 0.00

Mo 6.80

Si 93.20Ti, at.%

Te

mp

era

ture

, °C

L

L+Si

L+MoSi2

L+MoSi2+Si

L+TiSi2+Si L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+Si

TiSi2+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+MoSi2+Si

U, 1380°C

E, 1320°C

e8, 1330

e6, 1400

10

1200

1300

1400

1500

1600

1700

Ti 12.80

Mo 0.00

Si 87.20

Ti 0.00

Mo 6.80

Si 93.20Ti, at.%

Te

mp

era

ture

, °C

L

L+Si

L+MoSi2

L+MoSi2+Si

L+TiSi2+Si L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+Si

TiSi2+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+MoSi2+Si

U, 1380°C

E, 1320°C

e8, 1330

e6, 1400

Fig. 9a: Mo-Si-Ti.

Vertical section

through Si

isoconcentrate 80

mass%

402


MSIT®

Mo–Si–Ti

10

1200

1300

1400

1500

1600

1700

1800

Ti 16.40

Mo 0.00

Si 83.60

Ti 0.00

Mo 8.90

Si 91.10Ti, at.%

Te

mp

era

ture

, °C

L+Si

L+MoSi2+Si

L+TiSi2+Si

L+MoSi2

L

TiSi2+(Ti,Mo)Si2+Si

L+TiSi2

(Ti,Mo)Si2+MoSi2+Si

L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+SiE, 1320°C

U, 1380

e8, 1330

e6, 1400

10

1200

1300

1400

1500

1600

1700

1800

Ti 16.40

Mo 0.00

Si 83.60

Ti 0.00

Mo 8.90

Si 91.10Ti, at.%

Te

mp

era

ture

, °C

L+Si

L+MoSi2+Si

L+TiSi2+Si

L+MoSi2

L

TiSi2+(Ti,Mo)Si2+Si

L+TiSi2

(Ti,Mo)Si2+MoSi2+Si

L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+SiE, 1320°C

U, 1380

e8, 1330

e6, 1400

Fig. 9b: Mo-Si-Ti.

Vertical section

through Si

isoconcentrate 75

mass%

10

1200

1300

1400

1500

1600

1700

1800

1900

Ti 20.10

Mo 0.00

Si 79.90

Ti 0.00

Mo 11.10

Si 88.90Ti, at.%

Te

mp

era

ture

, °C

L

L+MoSi2

L+TiSi2 L+Si

L+TiSi2+Si

(Ti,Mo)Si2+MoSi2+Si

L+(Ti,Mo)Si2+Si

L+MoSi2+Si

(Ti,Mo)Si2+Si

TiSi2+(Ti,Mo)Si2+Si

E, 1320

U,

e8, 1330

e6, 1400

1380

10

1200

1300

1400

1500

1600

1700

1800

1900

Ti 20.10

Mo 0.00

Si 79.90

Ti 0.00

Mo 11.10

Si 88.90Ti, at.%

Te

mp

era

ture

, °C

L

L+MoSi2

L+TiSi2 L+Si

L+TiSi2+Si

(Ti,Mo)Si2+MoSi2+Si

L+(Ti,Mo)Si2+Si

L+MoSi2+Si

(Ti,Mo)Si2+Si

TiSi2+(Ti,Mo)Si2+Si

E, 1320

U,

e8, 1330

e6, 1400

1380

Fig. 9c: Mo-Si-Ti.

Vertical section

through Si

isoconcentrate 70

mass%

403


MSIT®

Mo–Si–Ti

20 10

1200

1300

1400

1500

1600

1700

1800

1900

Ti 24.00

Si 76.00

Mo 0.00

Ti 0.00

Si 86.40

Mo 13.60Ti, at.%

Te

mp

era

ture

, °C

L+MoSi2+Si

L

(Ti,Mo)Si2+Si

L+MoSi2

L+(Ti,Mo)Si2+Si

L+TiSi2

TiSi2+(Ti,Mo)Si2+Si

L+TiSi2+Si

(Ti,Mo)Si2+MoSi2+Si

1380

e6, 1400

e8, 1330

E, 1320

U

20 10

1200

1300

1400

1500

1600

1700

1800

1900

Ti 24.00

Si 76.00

Mo 0.00

Ti 0.00

Si 86.40

Mo 13.60Ti, at.%

Te

mp

era

ture

, °C

L+MoSi2+Si

L

(Ti,Mo)Si2+Si

L+MoSi2

L+(Ti,Mo)Si2+Si

L+TiSi2

TiSi2+(Ti,Mo)Si2+Si

L+TiSi2+Si

(Ti,Mo)Si2+MoSi2+Si

1380

e6, 1400

e8, 1330

E, 1320

U

Fig. 9d: Mo-Si-Ti.

Vertical section

through Si

isoconcentrate 65

mass% through the

invariant points E and

U13

20 10

1200

1300

1400

1500

1600

1700

1800

1900

2000

Ti 28.10

Mo 0.00

Si 71.90

Ti 0.00

Mo 16.30

Si 83.70Ti, at.%

Te

mp

era

ture

, °C

L

L+MoSi2

L+(Ti,Mo)Si2

L+TiSi2+(Ti,Mo)Si2

L+TiSi2+Si

L+TiSi2

L+(Ti,Mo)Si2+MoSi2

L+(MoSi2)+Si

L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+MoSi2+SiTiSi2+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+Sie, 1330

e6, 1400

20 10

1200

1300

1400

1500

1600

1700

1800

1900

2000

Ti 28.10

Mo 0.00

Si 71.90

Ti 0.00

Mo 16.30

Si 83.70Ti, at.%

Te

mp

era

ture

, °C

L

L+MoSi2

L+(Ti,Mo)Si2

L+TiSi2+(Ti,Mo)Si2

L+TiSi2+Si

L+TiSi2

L+(Ti,Mo)Si2+MoSi2

L+(MoSi2)+Si

L+(Ti,Mo)Si2+Si


(Ti,Mo)Si2+Sie, 1330

e6, 1400

Fig. 9e: Mo-Si-Ti.

Vertical section

through Si

isoconcentrate 60

mass%

404


MSIT®

Mo–Si–Ti

30 20 10

1200

1300

1400

1500

1600

1700

1800

1900

2000

Ti 32.40

Mo 0.00

Si 67.60

Ti 0.00

Mo 19.30

Si 80.70Ti, at.%

Te

mp

era

ture

, °C

L

L+(Ti,Mo)Si2+MoSi2L+MoSi2+Si


L+TiSi2+Si

L+(Ti,Mo)Si2

L+MoSi2

L+TiSi2

L+TiSi2+(Ti,Mo)Si2L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+Si

e6, 1400

E, 1320

U,1380e8, 1330

30 20 10

1200

1300

1400

1500

1600

1700

1800

1900

2000

Ti 32.40

Mo 0.00

Si 67.60

Ti 0.00

Mo 19.30

Si 80.70Ti, at.%

Te

mp

era

ture

, °C

L

L+(Ti,Mo)Si2+MoSi2L+MoSi2+Si


L+TiSi2+Si

L+(Ti,Mo)Si2

L+MoSi2

L+TiSi2

L+TiSi2+(Ti,Mo)Si2L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+Si

e6, 1400

E, 1320

U,1380e8, 1330

Fig. 9f: Mo-Si-Ti.

Vertical section

through Si

isoconcentrate 55

mass%

20 10

1200

1300

1400

1500

1600

1700

1800

1900

2000

Ti 27.57

Mo 5.76

Si 66.67

Ti 0.00

Mo 22.60

Si 77.40Ti, at.%

Te

mp

era

ture

, °C

L

L+(Ti,Mo)Si2+MoSi2

L+MoSi2+Si

TiSi2+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+MoSi2+Si

L+(Ti,Mo)Si2

L+TiSi2+(Ti,Mo)Si2

L+(Ti,Mo)Si2+Si

L+MoSi2

(Ti,Mo)Si2+Si

p7, 1510

U, 1380

E, 1320

e6, 1400

20 10

1200

1300

1400

1500

1600

1700

1800

1900

2000

Ti 27.57

Mo 5.76

Si 66.67

Ti 0.00

Mo 22.60

Si 77.40Ti, at.%

Te

mp

era

ture

, °C

L

L+(Ti,Mo)Si2+MoSi2

L+MoSi2+Si

TiSi2+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+MoSi2+Si

L+(Ti,Mo)Si2

L+TiSi2+(Ti,Mo)Si2

L+(Ti,Mo)Si2+Si

L+MoSi2

(Ti,Mo)Si2+Si

p7, 1510

U, 1380

E, 1320

e6, 1400

Fig. 9g: Mo-Si-Ti.

Vertical section

through Si

isoconcentrate 50

mass%

405


MSIT®

Mo–Si–Ti

10

1200

1300

1400

1500

1600

1700

1800

1900

2000

Ti 18.90

Mo 14.40

Si 66.70

Ti 0.00

Mo 26.40

Si 73.60Ti, at.%

Te

mp

era

ture

, °C

L

L+MoSi2+Si

L+(Ti,Mo)Si2+MoSi2

(Ti,Mo)Si2+MoSi2+Si

L+(Ti,Mo)Si2

L+MoSi2

L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+Si

e6, 1400

10

1200

1300

1400

1500

1600

1700

1800

1900

2000

Ti 18.90

Mo 14.40

Si 66.70

Ti 0.00

Mo 26.40

Si 73.60Ti, at.%

Te

mp

era

ture

, °C

L

L+MoSi2+Si

L+(Ti,Mo)Si2+MoSi2

(Ti,Mo)Si2+MoSi2+Si

L+(Ti,Mo)Si2

L+MoSi2

L+(Ti,Mo)Si2+Si

(Ti,Mo)Si2+Si

e6, 1400

Fig. 9h: Mo-Si-Ti.

Vertical section

through Si

isoconcentrate 45

mass%

406


MSIT®

N–Ni–Ti

Nitrogen – Nickel – Titanium

Honghui Xu, Yong Du, Zhaohui Yuan and Hailin Chen

Literature Data

A ternary nitride with an approximate formula of Ti0.7Ni0.3N was found to form on passing oxygen-free dry

ammonia over a finely powdered binary Ni-Ti alloy in the temperature range of 600 to 800°C [1954Sch].

A second phase, tentatively identified as Ti4Ni2N, was observed within the TiNi matrix after the Ti50Ni50

(at.%) alloys were melted in the presence of a controlled amount of nitrogen gas [1965Roz]. [1967Sto1,

1980Fuk, 1980Bon, 1981Mit, 1987Ego] investigated the TiN-Ni subsystem. [1987Ego] demonstrated that

the interaction of TiN with liquid nickel resulted in the N loss from TiN to form TiN0.51, (Ni), Ni3Ti, TiNi

and Ti2Ni, and the dissolution of Ni into the Ti nitride. [1980Fuk] reported a pseudobinary eutectic

occurring between (Ni) and stoichiometric TiN. The eutectic temperature was determined to be 1353 4°C

at a composition of about 11.4Ti-82.3Ni-6.3N (at.%). Further studies by the same researchers [1982Fuk]

were concerned with the mechanisms of grain growth and denitrification in Ti(C,N)-Ni sintered alloys. The

TiNi1-x+ NixTi1-x two-phase field was investigated by [1991Bin], and tentative phase relationships at

1100°C and pN2<105 Pa were presented. Approximately 10 sintered samples were prepared from nickel

(FSSS grain size 5 m, less than 200 ppm O and 50 ppm C), titanium (FSSS grain size 22.4-25.0 m,

maximum of 0.120% O and 0.024% N) and TiN (FSSS grain size 2-5 m, 0.1% C, 2.0% O) powders. The

samples were then annealed at 1100°C for 10 d. The water quenched samples were investigated by means

of X-ray diffraction (XRD), metallography, specific saturation magnetization and differential thermal

analysis (DTA). In order to ascertain the pseudobinary eutectic reported by [1979Fuk], the melting behavior

of TiN-Ni alloys was investigated by metallography and DTA. Metallographic investigation revealed no

eutectic structures, not even in those samples that had been prepared following [1979Fuk]. It was suggested

that the pseudobinary eutectic composition between TiN and Ni must be situated very close to the nickel

corner if the eutectic point reported by [1979Fuk] exists. Within the accuracy of the results, only

stoichiometric TiN was found to be in equilibrium with (Ni).

An isothermal section at 900°C, which is in agreement with previous work [1991Bin], was recently

established by [1998Lef] with the aid of thermodynamic modeling. No ternary compounds were observed

in the isothermal section. Approximately 60 samples were prepared by arc melting or high frequency

levitation melting of the high purity powder mixtures of Ti, Ni and TiN, which had been compacted in steel

dies without using lubricants. The investigation was based on XRD, metallography, scanning electron

microscopy (SEM), and electron probe microanalysis (EPMA) techniques. It was claimed that there is

convincing evidence for the two three-phase equilibria Ti2NiNx+Ti2N+TiNi1-x and Ti2NiNx+Ti2N+( Ti).

A significant nitrogen content of the ternary Ti2NiNx of up to about 11 at.% N at 900°C was revealed from

EMPA studies, which perhaps can account for the occurrence of Ti4Ni2N [1965Roz]. Despite extended

homogeneity regions existing for the phases Ti(N), Ti(N), TiN1-x, Ni(Ti) and TiNi, the solubilities of the

third components in these phases do not exceed about 0.5 at.%.

A thermodynamic assessment of the N-Ni-Ti system was performed by [1998Zen], who used the

experimental phase equilibrium data at 900°C from [1998Lef] and 1100°C from [1991Bin].

Binary Systems

The Ni-Ti and N-Ni binary phase diagrams are accepted from [Mas2]. The binary system N-Ti is essentially

due to the work of [1991Len] with some additional details concerning the phase boundary between ( Ti)

and Ti2N [1988Bar] as well as the precise location of the so-called ’ phase TiN0.51. The ’ phase was

found to form peritectoidally at 800 10°C via the reaction Ti2N+ TiN0.6 ’TiN0.51 [1991Etc]. However,

[1992Len] suggested that this phase is metastable. A recent critical assessment and thermodynamic

modelling of the N-Ti system (Fig. 1) has been performed by [1996Zen].

407


MSIT®

N–Ni–Ti

Solid Phases

Data for the solid phases are listed in Table 1. The stability of the ternary compound Ti0.7Ni0.3N postulated

by [1954Sch] need further study.

Isothermal Sections

Figures 2 to 4 show the thermodynamically calculated sections at 1800, 1500 and 1353°C with a pressure

of 5 Pa [1998Zen], respectively. Figure 5 exhibits the calculated isothermal section at 1 bar and at 1100°C

[1998Zen] where the computed phase equilibria agree with the experimental data [1991Bin].

Figure 6 presents the calculated isothermal section at 1 bar and 900°C [1998Zen] where the calculated phase

equilibrium relationships agree with those determined experimentally [1998Lef], but the Ti2N phase was

treated as stoichiometric. The calculated compositions of ( Ti), ( Ti) and Ti2Ni in the three phase equilibria

are obviously in contrast to the tentative experimental data, and the calculated homogeneity range of

TiN1-x is narrower than the experimentally determined one. It should be noted that the solid solubility of

N in Ti2Ni at 900°C was determined to be up to 11 1 at.% N. It seems that the reported phase “Ti4Ni2N”,

which was observed within the TiNi matrix after the Ti50Ni50 alloys were melted in the presence of a

controlled amount of nitrogen gas [1965Roz], is located within the N solubility of the Ti2Ni phase.


Figures 7 and 8 present the calculated vertical sections TiN - Ni, and TiN - 95Ni-5Ti (mass%) with a

pressure of 5 Pa, respectively, given by [1998Zen].

Thermodynamics

No experimental thermodynamic data are available.

A thermodynamic modeling of the ternary N-Ni-Ti system was performed by [1998Zen]. The liquid phase,

was described using a substitutional solution model. The ternary phases (Ni), ( Ti) and ( Ti) were

described by a two sublattice model (Ti,Ni)1(N,Va)b, in which Va denotes vacant interstitial sites on the

second sublattice, and with b = 1 for (Ni), b = 3 for ( Ti), and b = 0.5 for ( Ti), respectively. The solubility

of nitrogen in Ti2Ni was modeled using the sublattice formula (Ti)2(Ni)1(N,Va)0.5. Only the species Ti1,

Ni1, and N2 were included in the gaseous phase, which was treated as an ideal gas mixture in the calculation.

The thermodynamic description of the N-Ni-Ti system was applied to a study of the conditions for the

formation of nitride inclusions in nickel alloys containing titanium, the nitridation of Ni-Ti alloys, and the

denitrification, phase formation, and vaporization during the sintering of the TiN-Ni mixed powder

compacts.


Titanium nitride shows attractive properties: low density, high hardness, good electrical and thermal

conductivity, high melting point, and high wear and corrosion resistance. It can be used for optical and

abrasive layers, cutting tools, conductive layers, ballistic armor protections, and cermets. TiN particles are

also used as a strengthening phase in metal alloys and structural ceramics. A review of the solid-state

property of transition-metal nitrides was given by [1990Len]. [2000Bel] investigated the microstructure and

properties of titanium nitride-based composites. [1982Fuk] probed the mechanisms of grain growth in

Ti(C,N)-Ni sintered alloys.

The metallographic structures and hydrogenation characteristics of ternary alloys Ti4Ni2X (X = O, N, C)

were investigated by [2000Tak].

Miscellaneous

The effects of Ti on the solubility of N in dilute Ni alloys containing Ti were investigated through

thermodynamic calculation and experiment [1967Sto2], and the calculated results are in good agreement

with the experimental data. Experiments on the solubility of N in liquid Ni-Ti alloys at 1550°C and

408


MSIT®

N–Ni–Ti

p = 1atm demonstrated that Ti increases the solubility significantly only at concentrations of ~1%, and

subsequently with the formation of nitride when starting from a certain concentration. The thermodynamic

equation for the three-phase equilibrium of TiN, Ni and N2 was derived by [1967Sto1]. Calculations showed

that, the formation of nitrides in molten alloys (1550°C) is thermodynamically impossible if the content of

Ti is less than 1 %; and nitrides form at lower temperatures during cooling and solidification. Experiments

also showed that the minimum concentration of Ti to form nitrides in the melt (1550°C, 1atm) is 3-4%, and

this agrees with the calculated results.

An equivalent method for the calculation of the specific activity coefficients xN and x

H as a function of

temperature and concentration was given by [1987Sch]. By using this equivalent method, the partial

thermodynamic quantities, and the nitrogen and hydrogen solubility in nickel rich multicomponent alloys

were derived.

References

[1954Sch] Schoenberg, N., “The Tungsten Carbide and Nickel Arsenide Structures”, Acta Metall., 2,


[1965Roz] Rozner, A.G., Heintzelman, E.F., Buehler, W.J., Gilfrich, J.V., “Effect of Addition of

Oxygen, N, and H on Microstructure and Hardness of Cast TiNi Intermetallic Compound”,

Trans. Quart. ASM, 58, 415-418 (1965) (Crys. Structure, Experimental, 6)

[1967Sto1] Stomakhin, A.Y., Polyakov, A.Y., “Conditions for the Existence of a Nitride Phase in

Alloys of Ni with Ti, Zr and Al”(in Russian), Izv. VUZ, Chern. Metall., 10(3), 116-121

(1967) (Calculation, Experimental, 4)

[1967Sto2] Stomakhin, A.Y., Polyakov, A.Y., “Solubility of N and Formation of Nitrides in Molten

Alloys of Ni with Ti and Al”(in Russian), Izv. Akad. Nauk SSSR, Met., 4(2), 49-54 (1967)

(Calculation, Crys. Structure, Experimental, 5)

[1979Fuk] Fukuhura, M., Mitani, H., “On the Phase Relationship and the Dendritation During the

Sintering Process of the Titanium Nitride-Nickel Mixed Powder Compacts” (in Japanese),

Nippon Kinzoku Gakkaishi, 43(3), 169-174 (1979) (Experimental, Equi. Diagram, 4)

[1980Fuk] Fukuhura, M., Mitani, H., “The Phase Relationship and the Denitrification during the

Sintering Process of the Titanium Nitride-Nickel Mixed Powder Compacts”, Trans Jpn.

Inst. Met., 21(4), 211-218(1980) (Equi. Diagram, Experimental, 5)

[1980Bon] Bondar, V.T., “Reactions of Titanium Nitride in its Composites with Ni and Ni-Mo Alloy”,

Sov. Powder Metall. Met. Ceram., 19, 583-586 (1980), translated from Poroshk. Metall.,

(8), 85-90(1980) (Crys. Structure, Experimental, 11)

[1981Mit] Mitani, H., Nagai, H., Fukuhara, M., “ Denitrification of Titanium Nitride-Nickel Compacts

During Sintering”, Mod. Dev. Powder Metall, 14, 347-362(1981) (Experimental)

[1982Fuk] Fukuhara, M., Mitani, H., “Mechanisms of Grain Growth in Ti(C, N) Sintered Alloys”,

Powder Met., 25(2), 62-68 (1982) (Experimental, 20)

[1986Len1] Lengauer, W., Ettmayer, P., “The Crystal Structure of a New Phase in the

Titanium-Nitrogen System”, J. Less-Common Met., 120(1), 153-159 (1986) (Crys.



System”, J. Less-Common Met., 125, 127-134 (1986) (Crys. Structure, Experimental, 19)

[1987Sch] Schuermann, E., Sittard, M., Voelker, R., “Equivalent Influence of Alloying Elements on

Nitrogen and Hydrogen Solubility in Liquid Ternary and Multicomponent Nickel-Base

Alloys” (in German), Z. Metallkd., 78(6), 457-466 (1987) (Equi Diagram, Review,

Thermodyn., 53)

[1987Ego] Egorov, F.E., Smirnov, V.P., Timoifeeva, I.I., Mai, V.I., “Interaction of Titanium Nitride

with Liquid Nickel” (in Russian), Adgez. Rasplavov Paika Mater., 19, 59-63 (1987)

(Experimental, Kinetics, Mechan. Prop., 12)

[1988Bar] Bars, J.P., Etchessahar, E., Debuigne, J., “Metallurgy of the Ti-N System: Homogeneity,

Formation of Stable and Metastable Phases, Interphase Equilibria, Nitrogen Diffusion”,

409


MSIT®

N–Ni–Ti

Sixth World Conference on Titanium, Cannes, France, 1565-1570 (1988) (Crys. Structure,

Equi. Diagram, Experimental)

[1990Len] Lengauer, W., Ettmayer, P., “Recent Advances in the Field of Transition-Metal Refractory

Nitrides”, High Temp. -High Pressures, 22(1), 13-24 (1990) (Phys. Prop., Review, 42)

[1991Len] Lengauer, W., “The Titanium-Nitrogen System: A Study of Phase Reactions in the

Subnitride Region by Means of Diffusion Couples”, Acta Metall. Mater., 39(12),

2985-2996(1991) (Equi Diagram, Experimental, 21)

[1991Etc] Etchessahar, E., Sohn, Y.-U., Harmelin, H., Debuigne, J., “The Ti-N System: Kinetic,

Calorimetric, Structure and Metallurgical Investigation of the ’-TiN0.51 Phase”,

J. Less-Common Met., 167, 261-281(1991) (Equi. Diagram, Crys. Structure,

Experimental, 9)

[1991Bin] Binder, S., Lengauer, W., Ettmayer, P., “The Ti-N-Ni System: Investigations Relevant for

Cerment Sintering”, J. Alloys Compd., 177(1), 119-127 (1991) (Equi. Diagram,

Experimental, #, 18) see also: [1994Mch] McHale, A.E., “XVIII. Nitrogen Plus Two

Metals”, Phase Equilibria Diagrams, Phase Diagrams for Ceramists, 10, 450 (1994) (Equi.


[1992Len] Lengauer, W., “A Study of ’-TiN1-x Formation in Temperature Gradient Diffusion

Couples”, J. Alloys Compd., 289-297 (1992) (Equi. Diagram, Experimental, 15)

[1996Zen] Zeng, K., Schmid-Fetzer, R., “Critical Assessment and Thermodynamic Modelling of the

N-Ti System”, Z. Metallkd., 87(7), 540-554 (1996) (Calculation, Equi. Diagram,

Review, 69)

[1998Zen] Zeng, K., Schmid-Fetzer, R., Rogl, P., “ A Thermodynamic Analysis of Cermet Sintering

of TiN-Ni Powder Mixtures”, J. Phase Equilib., 19(2), 124-135 (1998) (Calculation, Equi.


[1998Lef] Le Friec, Y., Rogl, P., Bauer, J., Bohn, M., Zeng, K., Schmid-Fetzer, R., “Investigation of

the Nitrogen-Nickel-Titanium System: the Isothermal Section at 900°C”, J. Phase Equilib.,

19(2), 112-123 (1998) (Crys. Structure, Equi. Diagram, Experimental, #, *, 24)

[2000Bel] Bellosi, A., Monteverde, F., “Microstructure and Properties of Titanium Nitride and

Titanium Diboride-Based Composites”, Key Eng. Mater., 175-176, 139-148 (2000)


[2000Tak] Takeshita, H.T., Tanaka, H., Kuriyama, N., Sakai, T., Uehara, I., Haruta, M.,

“Hydrogenization Characteristics of Ternary Alloys Containing Ti4Ni2X (X = O, N, C)”,

J. Alloys Compd., 311, 188-193 (2000) (Experimental, Phys. Prop., 7)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 [Mas2]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06 [Mas2]

(Ni)

< 1445

Ni1-xTix

cF4

Fm3m

Cu

a = 352.40

a = 356.1

at x = 0 [Mas2]

at x = 0.09 [V-C2]

410


MSIT®

N–Ni–Ti

( N)

-237 - 210

hP4

P63/mmc

N

a = 405.0

c = 660.4

[Mas2]

( N)

<- 237

cP8

Pa3

N

a = 566.1 [Mas2]

Ti2N

< 1080

tP6

P42/mnm

anti-TiO2

a = 494.52

c = 303.42

31-33 at.% N

[V-C2]

TiN1-x

< 3290

cF8

Fm3m

NaCl

a = 424.42

a = 420.50

at x = 0 [1992Len]

at x = 0.61

[V-C2]

Ti4N3-x

1291 - 1078

hR8

R3m

V4C3

a = 297.95

c = 2896.49

31.5 at.% N [1986Len1]

Ti3N2-x

1103 - 1066

hR6

R3m

VTa2C2

a = 298.09

c = 2166.42

29 at.% N [1986Len2]

Ti2Ni

< 984

cF96

Fd3m

Ti2Ni

a = 1127.8

a = 1131.93

Dissolves N. Denoted as Ti2Ni(N)x

[Mas2]

[V-C2]

TiNi

1310 - 630

Ti1-xNix

cP2

Pm3m

CsCl

a = 301.2

a = 298.6 x = 0.55 [Mas2]

x = 0.45 [V-C2]

TiNi3< 1380

hP16

P63/mmc

TiNi3

a =510.88

c = 831.87

[Mas2]

[V-C2]

Ni3N

600

hP8

P6322

ReO3

a = 461.6

c = 429.8

[Mas2]

[V-C2]

Ni4N

250

cP5

Pm3m

CaO3Ti

a = 374.0 [Mas2]

[V-C2]

NiN6 Unknown [Mas2]

* , Ti0.7Ni0.3N hP2

P6m2

WC

a = 294

c = 289

[1954Sch]

Ti0.7Ni0.3N was prepared at various

temperatures between 600 and 800°C

[1954Sch]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

411


MSIT®

N–Ni–Ti

10 20 30 40

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

Ti Ti 50.00

N 50.00N, at.%

Te

mp

era

ture

, °C

3060

L+gas (1 bar)

δTiN1-x

2347

1995

1278

1104

10631075

1082

Ti3N2

Ti4N3

εTi2N

(βTi)1670°C

882°C

(αTi)

L

Fig. 1: N-Ni-Ti.

N-Ti phase diagram

[1996Zen]

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni

N Data / Grid: at.%

Axes: at.%

δ+L+gas

(αTi)+L+δ

(αTi)

L

δ+gas

(αTi)+L

δ

(βTi)gas

Fig. 2: N-Ni-Ti.


section at 1800°C

[1998Zen]

412


MSIT®

N–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni

N Data / Grid: at.%

Axes: at.%

L+gas+δ

δ+(αTi)+L

(αTi)

(βTi)

δ+L

(αTi)+L

gas+L

L

δ

gas

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni

N Data / Grid: at.%

Axes: at.%

L+gas+δ

δ+(αTi)+L

(αTi)

(βTi)

TiNi3

(Ni)

δ+L

(αTi)+L

L

δ+TiNi3+L

L

Fig. 3: N-Ni-Ti.


section at 1500°C

[1998Zen]

Fig. 4: N-Ni-Ti.


section at 1353°C

[1998Zen]

413


MSIT®

N–Ni–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni

N Data / Grid: at.%

Axes: at.%

(Ni)+δ+gas(1bar)

δTiN1-x

εTi2N

(αTi)

(βTi)Ti

2Ni TiNi TiNi

3(Ni)

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni

N Data / Grid: at.%

Axes: at.%

(Ni)+δ+gas (1bar)δTiN

1-x

ζTi4N

3

(αTi)

(βTi)TiNi TiNi

3(Ni)L

ηTi3N

2

Fig. 5: N-Ni-Ti.


section at 1100°C

[1998Zen]

Fig. 6: N-Ni-Ti.


section at 900°C

[1998Zen]

414


MSIT®

N–Ni–Ti

20 40 60 80

750

1000

1250

1500

1750

2000

Ti 50.01

Ni 0.00

N 49.99

Ni

Ni, at.%

Te

mp

era

ture

, °C

gas (5 Pa)

gas+δ

L+gas+δ

(Ni)+gas+δ

(Ni)+gas

L+gas

70 80 90

1200

1300

1400

1500

1600

Ti 17.45

Ni 69.45

N 13.10

Ti 6.06

Ni 93.94

N 0.00Ni, at.%

Te

mp

era

ture

, °C

Z(Ni)+gas

L+gas

L+gas+δ

(Ni)+gas+δ

Fig. 7: N-Ni-Ti.

Calculated vertical

section TiN - Ni at 5

Pa [1998Zen]

Fig. 8: N-Ni-Ti.

Calculated vertical

section from TiN to

95Ni-5Ti (mass%) at

5 Pa [1998Zen]

415


MSIT®

N–Ti–V

Nitrogen – Titanium – Vanadium

Zhaohui Yuan, Yong Du, Honghui Xu, Wei Xiong

Literature Data

The ternary N-Ti-V system was investigated by [1950Duw, 1966Str, 1972Kie, 1973Shu, 1974Bar,

1993Gui, 1996Eno, 1998Gui, 2002Kro]. The literature data up to 1989 were reviewed by [1996Eno].

[1950Duw, 1966Str, 1972Kie] reported a mutual solubility of TiN and VN and determined the lattice

parameters along the TiN-VN section. The isothermal section at 1200°C at compositions of less than

25 at.% N and the V-TiN pseudobinary section were constructed by [1973Shu] using metallographic

observation, X-ray diffraction (XRD), and differential thermal analysis (DTA). [1974Bar] measured the

solubility of titanium nitride in V at 1200, 1500, and 1800°C. [1973Shu, 1974Bar] prepared alloys in an arc

melting furnace under a mixture of Ar and high purity N. The alloys were then annealed at 1200, 1500 and

1800°C for a long time. [1993Gui] investigated the phase equilibria in the N-Ti-V system at 1000 and

700°C. The alloys used by [1993Gui] were prepared by arc melting commercial powders of the pure

elements and nitrides under an Ar atmosphere followed by annealing at 700 and 1000°C for 1400 and 400 h,

respectively. The phases were identified by means of XRD and metallographic observation, and the phase

compositions were measured by electron probe microanalysis (EPMA). The solubility of V in the Ti2N

phase was < 2 at.% at 1000°C; in Ti it was between 1 and 2 at.% at 1000°C and < 2 at.% at 700°C. The

maximum solubility of Ti in V2N1-y was 5 at.% at 1000°C.

[1998Zen] assessed the thermodynamic functions of the phases in the N-Ti-V system and calculated the

phase equilibria at 1000, 1200 and 1800°C. The isothermal sections calculated at 1000 and 1200°C are in

reasonable agreement with the experimental data of [1993Gui] and [1974Bar], respectively. However, the

experimentally determined solubility of N and Ti in V at 1200 to 1800°C according to [1974Bar] is much

less than that calculated by [1998Zen]. Also, [1998Zen] showed that the V-TiN system is not pseudobinary

with a eutectic reaction of L + as claimed by [1973Shu]. From the calculation, [1998Zen] suggested that

there is a transitional invariant reaction L+ ´ + at 1697°C in the V-TiN vertical section. This

temperature is lower than the invariant temperature suggested by [1973Shu] by 170°C.

[1998Gui] established diffusion paths in the isothermal section at 1200°C by investigating the nitridation of

Ti-V diffusion couples. After diffusion annealing, the couples were cut perpendicularly to the diffusion

interface and polished. The concentration profiles were measured using EPMA. The resulting EPMA

determinations were accurate to within 0.5 at.% for Ti and V and 1 at.% for N. The experimental results of

[1998Gui] confirmed the calculated isothermal section at 1200oC given by [1998Zen].

A series of Ti rich N-Ti-V alloys containing about 20 at.% V and between 11 and 15 at.% N have been

studied experimentally and an attempt has been made to calculate the isothermal section at 1200°C

[2002Kro]. Ingots were prepared by melting high purity Ti and V in an arc melting furnace. The gas

atmosphere consisted of mixtures of Ar and N in three different ratios, flowing at a constant rate with a

0.021 MPa overpressure. Specimens were annealed at 1200°C for 1 to 3 h in sealed evacuated quartz

capsules. In order to have a complete understanding of the nature of phase transformations and phase

equilibria, water-quenched specimens were analyzed using XRD, scanning electron microscopy (SEM),

back scattered electron microscopy (BSE), transmission electron microscopy (TEM), energy dispersive

X-ray spectrometer (EDS), parallel electron energy loss spectrometry (PEELS), and wavelength dispersive

X-ray analysis (WDX). The agreement between the experimental results and the calculations of [2002Kro]

and [1998Zen] is good.

Binary Systems

The N-V and Ti-V binary systems are accepted from [1997Du] and [1998Sau], respectively. The N-Ti

binary phase diagram was calculated by [1996Zen] without consideration of the gas phase and by [1998Zen]

who subsequently included a description of gas phase. In the present work, the N-Ti phase diagram is

416


MSIT®

N–Ti–V

accepted from [1998Zen]. The binary N-Ti, N-V and Ti-V phase diagrams are shown in Figs. 1 to 3,

respectively.

Solid Phases

Table 1 lists the crystal structure data for the solid phases in the ternary N-Ti-V system. No ternary

compound has been observed experimentally. According to the experimental data of [1991Len], the binary

phases , and ,Ti2N have narrow homogeneity ranges (about 1 at.% N for and and about 2 at.% for ).

The phases TiN and VN form a complete series of solid solutions.


Table 2 shows the calculated invariant equilibria according to [1998Zen]. A partial reaction scheme is

shown in Fig. 4.

Liquidus Surface

Figure 5 presents the calculated liquidus surface according to the thermodynamic modeling of [1998Zen].

Isothermal Sections

The calculated [1998Zen] isothermal sections at 1800, 1200, and 1000°C are shown in Figs. 6 to 8,

respectively. These computed sections agree well with the limited experimental data [1972Kie, 1973Shu,

1993Gui, 1998Gui, 2002Kro].


The calculated V-TiN vertical section of the N-Ti-V system is presented in Fig. 9 according to [1998Zen].

Thermodynamics

There are no experimental thermodynamic data available for the ternary N-Ti-V system. [1998Zen]

modeled the N-Ti-V system by consideration of the experimental data available in the literature.


[1982Ban] investigated the structure and hardness of alloys along the V-TiN isopleth at high temperature.

A maximum hardness in excess of 2000 MPa was reached with a content of the strengthening phase between

3 and 6 vol.%.

Nanostructured monolayer and multilayer TiN/VN films are of primary importance in modern science and

technology because of their high hardness, high bulk modulus and shear modulus [2001And]. Mixed Ti and

V nitrides have been developed as coatings for cutting tools. In addition, fine particles of Ti and V nitrides

in steels have been used for both grain size control and precipitation strengthening.

Miscellaneous

[1978Ven] described the structure and stability of the transition metal nitrides in different ternary systems

including N-Ti-V.

[1998Bur] investigated structural changes in Ti alloys with a V content of 25 at.% and N contents of 11.5,

13.3 to 16.5 at.% annealed at 1000°C. Increasing the N content leads to the appearance of twinning in the

hcp structure.

[1999Aga] investigated the influence of H on the structural change of the fcc phase along the TiN-VN

section.

417


MSIT®

N–Ti–V

References

[1950Duw] Duwez, P., Odell, F., “Phase Relationships in the Binary Systems of Nitrides and Carbides

of Zirconium, Niobium, Titanium, and Vanadium”, J. Electrochem. Soc., 97, 299-304


[1966Str] Strashinakaya, L.V., Mironova, N.V., “Contact Interaction of TiN, Zr, and V in Vacuum”

(in Russian), Izv. Akad. Nauk SSSR, Met., 4, 143-146 (1966) (Experimental, Equi.

Diagram, 6)

[1972Kie] Kieffer, R., Nowotny, H., Ettmayer, P., Dufek, G., “Miscibility of the Nitrides and Carbides

of Transition Metals” (in German), Metall. Technik, 26, 701-708 (1972) (Experimental,


[1973Shu] Shurin, A.K., Barabash, O.M., “Phase Equilibria in Alloys of V with Ti, Zr, and Hf Nitride”

(in Russian), Akad. Nauk Ukr. SSR, Metallofiz., 45, 84-87 (1973) (Experimental, Equi.

Diagram, 11)

[1974Bar] Barabash, O.M., Shurin, A.K., “Solubility of TiN, ZrN and HfN in V”(in Russian), Izv.

Akad. Nauk SSSR, Met., 4, 194-197 (1974) (Experimental, Equi. Diagram, 4)

[1978Ven] Vendl, A., “About Nitride Chemistry of Transition Metals” (in German), Planseeber.

Pulvermetall., 26, 233-243 (1978) (Crys. Structure, Equi. Diagram, Review, 29)

[1982Ban] Bankovskiy, O.I., Barabash, O.N., Moiseev, V.F., “The Lasting Hardness of the Alloys in

V-TiN, V-TiN, V-ZrN and V-HfN Systems” (in Russian), Poroshk. Metall., (4), 95-99


[1986Len1] Lengauer, W., Ettmayer, P., “The Crystal Structure of a New Phase in the

Titanium-Nitrogen System”, J. Less-Common Met., 120, 153-159 (1986) (Crys. Structure,

Experimental, 24)


System”, J. Less-Common Met., 125, 127-134 (1986) (Crys. Structure, Experimental, 19)

[1991Len] Lengauer, W., “The Titanium-Nitrogen System: A Study of Phase Reactions in the

Subnitride Region by Means of Diffusion Couples”, Acta Metall. Mater., 39(12),


[1993Gui] Guillou, A., Bauer, J., Debuigne, J., “On the Ternary System Ti-N-V”, Titanium 92: Sci

Technol., Proc. Symp., 1992, Warrendale, PA, TMS, USA., 1, 747-754 (1993)


[1996Eno] Enomoto, M., “The N-Ti-V System”, J. Phase Equilib., 17(3), 248-252 (1996)

(Assessment, Equi. Diagram, 10)

[1996Zen] Zeng, K., Schmid-Fetzer, R., “Critical Assessment and Thermodynamic Modelling of the

N-Ti System”, Z. Metallkd., 87(7), 540-554 (1996) (Calculatuon, Equi. Diagram,

Review, #, 69)

[1997Du] Du, Y., Schmid-Fetzer, R., Ohtani, O., “Thermodynamic Assessment of the V-N System”,

Z. Metallkd., 88, 545-556 (1997) (Equi. Diagram, Thermodyn., Assessment, #, 53)

[1998Bur] Bursik, J., Chen, J.K., Kroupa, A., Weatherly, G.C., “Microstructural Changes of the HCP

Phase in Ti-25V-N Alloys Annealed at 1273K”, Scr. Mater., 39(8), 1139-1144 (1998)


[1998Gui] Guillou, A., Ansel, D., Debuigne, J., “Nitridation of Titanium Vanadium Diffusion

Couples”, Scr. Mater., 38(6), 981-989 (1998) (Equi. Diagram, Experimental, 16)

[1998Sau] Saunders, N., “System Ti-V”, in “COST 507 Thermochemical Database for Light Alloys”,

Ansara, I., Dinsdale, T., Rand, M.H. (Eds.), European Communities, Luxemburg, Vol. 2,

297-298 (1998) (Assessment, Equi. Diagram, #, 1)

[1998Zen] Zeng, K., Schmid-Fetzer, R., “Thermodynamic Assessment and Applications of the Ti-V-N

System”, Mater. Sci. Technol., 14, 1083-1091 (1998) (Assessment, Equi. Diagram,

Thermodyn., #, *, 34)

[1999Aga] Agadzhanyan, N.N., Akopyan, A.G., Beibutyan,V.M., Ter-Galstyan, O.P.,

Dolukhanyan, S.K., Shekhtman, V.S., “Complex Hydrides and Hydronitrides of Transition

418


MSIT®

N–Ti–V

Metals Prepared by SHS. II. Systems Involving Group V Metals and Nitrogen (Ti-V-N-H,

Ti-Nb-N-H-, Zs-Nb-N-H)”, Powder Metall. Met: Cer., 38(3-4), 176-178 (1999) (Crys.


[2001And] Andrievski, R.A., “Superhard Materials Based on Nanostructured High-Melting Point

Compounds: Achievements and Perspectives”, Iner. J. Ref. Met. Hard Mater., 19(4-6),

447-452 (2001) (Mechan. Prop., Review, 59)

[2002Kro] Kroupa, A., Bursik, J., Svoboda, M., Chen, J.K., Weatherly, G.C., “Phase Transformations

and Phase Equilibria in Ti-25V-N System at 1200°C”, Mater. Sci. Technol., 18, 13-20

(2002) (Calculation, Experimental, Equi. Diagram, 14)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( TixV1-x)

( Ti)

1670 - 882

(V)

< 1910

cI2

Im3m

W

a = 330.65

a = 302.4

at 0 x 1 [1998Sau],

dissolves about 3 at.% N at 1200°C

[1998Zen]

at x = 1 [Mas2],

dissolves up to 6.8 at.% N at 1995°C

[1996Zen]

at x = 0 [Mas2],

dissolves up to 17 at.% N at 1859°C

[1997Du]

( Ti)

< 882

hP2

P63/mmc

Mg a = 295.06

c = 468.35

dissolves up to 21.6 at.% N at 1278°C

[1996Zen]

at 25°C [Mas2]

( N)

< -253

tP4

P42/mnm

N

a = 395.7

c = 510.9

at pressure > 3.3 GPa [Mas2]

( N)

-237.54 to -210.004

hP4

P63/mmc

N

a = 405.0

c = 660.4

[Mas2]

( N)

< -237.54

cP8

Pa3

N

a = 566.1 [Mas2]

1104 - 1063

hR6

R3m

VTa2C2

a = 298.09

c = 2166.42

[1996Zen]

at 29 at.% N [1986Len2]

1278 - 1075

hR8

R3m

V4C3

a = 297.95

c = 2896.5

[1996Zen]

at 31.5 at.% N [1986Len1]

, Ti2N

< 1082

tP6

P42/mnm

TiO2

a = 494.52

c = 303.42

[1996Zen]

[V-C2]

419


MSIT®

N–Ti–V


, TixV1-xN1-y-z

, TixN1-y-z

< 3060

, V1-xN1-y-z

< 2119

cF8

Fm3m

NaCl

a = 421.6

a = 419.2

a = 416.3

a = 424.42

a = 413.47

at 0 x 1,

0 y 0.35, 0 z 0.5 and 1000 to

1200°C [1998Zen]

at x =0.668, y = 0, z = 0

(33.4Ti-16.6V (at.%)) [1972Kie]

at x = 0.5, y = 0, z = 0

(25Ti-25V (at.%)) [1972Kie]

at x = 0.332, y = 0, z = 0

(16.6Ti-33.4V (at.%)) [1972Kie]

[1996Zen]

at x = 1, y = 0, z = 0 (50 at.% N) [V-C2]

[1997Du]

at x = 0, y = 0, z = 0 (50 at.% N) [V-C2]

´, V2N1-y

< 2409

hP9

P31m

Fe2N

a = 491.7 0.3

c = 456.8 0.3

[1997Du]

[V-C2]


N Ti V

L + gas ´ + 2378 U1 L

gas

´

25.72

99.96

31.30

41.35

9.25

0.00

6.38

20.20

65.03

0.04

62.32

38.45

L + ´ + 1697 U2 L

´

11.02

27.82

9.93

42.32

12.51

15.95

5.78

50.82

76.47

56.23

84.29

6.86

L + ( Ti) + 1571 U3 L

( Ti)

5.44

21.26

4.40

31.81

49.77

73.64

52.89

66.01

44.79

5.1

42.71

2.18

L + 1560 emin L

5.75

4.04

33.41

42.81

41.63

64.21

51.44

54.33

2.38

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

420


MSIT®

N–Ti–V

90 80 70 60

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

Ti Ti 50.00

N 50.00Ti, at.%

Te

mp

era

ture

, °C

3060

L+gas (1 bar)

δ

2347

1995

1278

1104

1063 1075

1082

η

ζ

ε

(βTi)1670°C

882°C

(αTi)

L

Fig. 1: N-Ti-V.

N-Ti phase diagram at

1 bar [1998Zen]

10 20 30 40

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

V Ti 0.00

V 50.00

N 50.00N, at.%

Te

mp

era

ture

, °C

L+gas

L

(V)

1910°C 1859°C

2409°C

2119°C

α´ δ

α´+δ(V)+α´

L+α´

Fig. 2: N-Ti-V.

N-V phase diagram at

1 bar [1997Du]

421


MSIT®

N–Ti–V

20 40 60 80

0

250

500

750

1000

1250

1500

1750

2000

Ti V

V, at.%

Te

mp

era

ture

, °C

L

β

(αTi)

1670°C

882°C

1608°C

1910°C

Fig. 4: N-Ti-V. Partial reaction scheme

N-V A-B-C

l + gas α´

2409 p2

N-Ti-V

L + gas α´ + δ2378 U1

Ti-N

l + gas δ3060 p

1

l + α´ β1859 p

5

l + δ (αTi)

2347 p3

l + (αTi) β1995 p

4

L + α´ β + δ1697 U2

L + (αTi) β + δ1571 U3

gas+α´+δ

L+α´+δ

L β + δ1560 min

α´+β+δL+β+δ

L+β+δ (αTi)+β+δ

β+δ

Fig. 3: N-Ti-V.

Ti-V phase diagram

[1998Sau]

422


MSIT®

N–Ti–V

20

40

60

80

20 40 60 80

20

40

60

80

Ti V

N Data / Grid: at.%

Axes: at.%

gas+δ

δ

L+α´+δ

L

α´

β

(βTi)

(αTi)

L+β+α´L+δ+(αTi)

20

40

60

80

20 40 60 80

20

40

60

80

Ti V

N Data / Grid: at.%

Axes: at.%

U1

U2

α´δ

β

U3

(αTi)

gas

emin

L+gas+δ

Fig. 6: N-Ti-V.



pN2 = 8.106 bar

[1998Zen]

Fig. 5: N-Ti-V.

Calculated liquidus

surfaces at pN2 = 1bar

[1998Zen]

423


MSIT®

N–Ti–V

20

40

60

80

20 40 60 80

20

40

60

80

Ti V

N Data / Grid: at.%

Axes: at.%

α´

β

δ

ζ

(αTi)

gas+δ

β+δ+(αTi)

β+α´+δ

20

40

60

80

20 40 60 80

20

40

60

80

Ti V

N Data / Grid: at.%

Axes: at.%

(αTi)

δ δ

α´

β

gas+δ

β+δ+α´

β+ε+δ

β+ε+(αTi)

ε

Fig. 7: N-Ti-V.



pN2 = 1 bar [1998Zen]

Fig. 8: N-Ti-V.



pN2 = 8.106 bar

[1998Zen]

424


MSIT®

N–Ti–V

20 40 60 80

0

250

500

750

1000

1250

1500

1750

2000

2250

2500

2750

3000

3250

3500

Ti 50.00

V 0.00

N 50.00

V

V, at.%

Te

mp

era

ture

, °C

1697°C

L

β+α´+δ β+δ

β+δ+ε β+ε

β

L+α´

L+βL+α´+β

L+δ+α´

L+δ

L+gas

δ+gas

δ

δ+α´

δ+L+gas

L+β+δ

47.517.3

59

74.5

82.6

78.6

Fig. 9: N-Ti-V.

Calculated vertical

section TiN-V

at pN2 = 1 bar

[1998Zen]

425


MSIT®

Ni–Pd–Ti

Nickel – Palladium – Titanium

Gautam Ghosh

Literature Data

There are five reports on the ternary phase equilibria studies [1962Rhy, 1980Bor, 1981Bor, 1982Bor,

1992Ger]; however, these were restricted to only part of the ternary system. In an effort to search for suitable

brazing alloys, [1962Rhy] reported a vertical section of Pd-rich compositions with 10 mass% Ti. They used

metallography and thermal analysis techniques. [1980Bor, 1981Bor] determined isothermal sections at 700

and 900°C for the composition range Ti-TiPd-NiTi, and two vertical sections NiTi-TiPd and NiTi2-PdTi2.

Ternary alloys were prepared by inductive levitation using iodide grade Ti, metallic Pd and electrolytic Ni,

but the purities were not mentioned. The alloys were annealed at 700 and 900°C for up to 300 h. X-ray

diffraction, differential thermal analysis, and metallography techniques were employed to determine the

phase equilibria. [1982Bor] determined a partial isothermal section at 400°C for the composition range

Ti-TiPd-NiTi. [1992Ger] reported an isothermal section at 800°C representing the composition range

Ni-Ni3Ti-Pd3Ti-Pd. To search for suitable alloys as filler material for joining Ti alloys, [1994Koe]

investigated thermophysical and thermochemical properties of Ti-(5 to 35)Ni-(5 to 10) Pd (mass%) alloys

which were prepared using 99.99% Ni, 99.95% Pd and 99.99% Ti. The alloys were prepared by arc-melting

in an argon atmosphere. The alloys were characterized using X-ray diffraction, metallography and thermal

analysis techniques. They reported DTA thermograms, selected microstructures and specific heat of the

alloys.

Binary Systems

The Ni-Pd, Ni-Ti and Pd-Ti binary phase diagrams are accepted from [1984Nas], [2004Ted] and

[1982Mur], respectively.

Solid Phases

TiNi and TiPd form a continuous solid solution [1980Bor]. NiTi2 dissolves about 5 at.% Pd, and Ti2Pd

dissolves about 7 at.% Ni [1980Bor, 1981Bor]. In Ti2Ni and Ti2Pd, Pd and Ni reside primarily on the Ni

and Pd sublattices, respectively. TiNi3 and TiPd3 form a continuous solid solution at 800°C [1982Bor].

Substitution of Ni by Pd in TiNi3 causes a rapid increase in both a and c lattice parameters [1982Bor].

There is no ternary phase in this system. The details of the crystal structures and lattice parameters of the

solid phases are listed in Table 1.

Solid – Solid Displacive Phase Transformations

Binary B2-NiTi undergoes martensitic transformation at low temperature to a monoclinic structure,

commonly known as the B19’ martensite [1992Shi]. However, at slightly richer Ni compositions, the

presence of dislocations, or the presence of metastable Ti3Ni4 precipitates are known to promote another

displacive transformation to a rhombohedral structure preceding B19’ transformation, commonly known as

the R phase [1997Har]. Depending on the composition and thermal history of binary TiNi, and with the

addition of Pd, the transformation temperatures B2 R B19’ may be well separated. The presence of

R-phase is very useful for shape memory applications which rely on small thermal hysteresis.

Binary B2-PdTi undergoes martensitic transformation at low temperature to an orthorhombic structure,

commonly known as the B19 martensite [1982Mat, 1983Siv, 1985Mat, 1993Mat, 1997Sem]. Figure 1

shows the structural diagram for displacive transformations along NiTi-PdTi section [1985Mat]. As seen,

depending on the alloy composition the transformation may proceed as B2 R B19’, B2 B19 B19’

and B2 B19. In Fig. 1, B2( ) represents an incommensurate phase [1982Hwa] where there is a lattice

displacement wave along 2/3<111> <111>B2 [1985Mat]. The B2 B2( ) is believed to be either a weakly

first-order or a second-order phase transformation. The stability of R phase increases by adding up to about

426


MSIT®

Ni–Pd–Ti

3 at.% Pd in NITi. The B2 B19 transformation temperature decreases by adding Ni to PdTi. The structural

details of B19, B19’ and R phases, and also composition dependence of their lattice parameters are given

in Table 1.


[1980Bor, 1981Bor] established the pseudobinary section NiTi-TiPd, which is shown in Fig. 2.

Isothermal Sections

Figures 3, 4, 5 and 6 show partial isothermal section at 900°C [1980Bor, 1981Bor], 800°C [1992Ger],

700°C [1980Bor, 1981Bor] and 400°C [1982Bor], respectively. Adjustments have been made to comply

with the accepted binary phase diagrams. For example, in this work NiTi2 is considered as a stoichiometric

compound [2004Ted]. One discrepancy was noted between the accepted Pd-Ti phase diagram and the

partial isothermal section at 800°C [1992Ger]. Not only TiPd4 intermetallic is shown to be present at 800°C,

it is also reported to dissolve up to 25 at.% Ni. However, the existence of TiPd4 in the binary Pd-Ti system

has not been confirmed [1982Mur]. Therefore, solid solution (Pd,Ni)4Ti has not been included in Fig. 4.


Figure 7 shows an isopleth at a constant Ti content of 66.7 at.% [1980Bor, 1981Bor]. Figure 8 shows the

temperature-composition section along the line of constant Pd:Ni mass ratio of 3:2 [1962Rhy]. The Ni-Pd

alloy with composition 3:2 (mass ratio) is characterized by a congruent melting at 1237°C.

Thermodynamics

[1994Koe] measured specific heat of Ti-(5 to 35)Ni-(5 to 10)Pd (mass%) alloys, both in as-cast and

annealed (at 800°C for 12 h) states. However, these alloys were not single phase. Besides ( Ti), the alloys

also contain Ti2Ni and Ti4Pd phases.


B2-(NiPd)Ti alloys are promising high-temperature shape memory alloys. As a result, a large number of

studies have been carried out to characterize their transformation and shape memory behaviors, e.g.

[1982Mat, 1983Siv, 1985Mat, 1993Mat, 2000Shi, 2000Xu, 2001Tia, 2002Tia], and also oxidation behavior

[2003Tia]. [2000Shi] observed (Ni,Pd)4Ti3 precipitates in aged (at 400°C for 10 h) Pd and Ni rich alloys of

Ti48NixPd52-x (x=0 to 52); however, the structure of (Ni,Pd)4Ti3 precipitates was reported to be different

from the metastable Ti3Ni4 precipitates in binary TiNi alloys. The potential and problems of (NiPd)Ti alloys

for high-temperature shape memory applications have been discussed by [1999Ots].

Miscellaneous

In as-cast alloys of Ti-(5 to 35)Ni-(5 to 10)Pd (mass%), [1994Koe] observed a metastable phase Ti3Ni

which disappears after annealing at 800°C for 12 h. Thermal analysis of these alloys shows a first melting

peak between 900 and 960°C, and alloys containing more than 25 mass% Ni show a second melting peak

around 1070°C.

References

[1962Rhy] Rhys, D.W., Berry, R.D., “The Development of Pd Brazing Alloys”, Metallurgia, 66,

255-263 (1962) (Experimental, Equi. Diagram, #, *, 5)

[1980Bor] Boriskina, N.G., Kenina, E.M., “Phase Equilibria in the Ti-TiPd-TiNi System Alloys”,

Titanium 80, Science and Technology, Proc. 4th International Conference on Titanium,

Kimura, H., Izumi, O., (Ed.), Kyoto, Vol. 4, 2917-2927 (1980) (Equi. Diagram,


427


MSIT®

Ni–Pd–Ti

[1981Bor] Boriskina, N.G., Kenina, E.M., “Phase Structural of Alloys of Ti-TiPd-TiNi System at

Temperatratures 700 and 900°C” (in Russian), in “Phase Equilibria in Metallic Alloys”,

Nauka, Moscow, 131-135 (1981) (Crys. Structure, Experimental, Equi. Diagram, Review,

#, *, 10)

[1982Bor] Boriskina, N.G., Kenina, E.M., Tumanova, T.A, “Phase Equilibrium and Some Properties

of Alloys of the Ti-TiPd-TiNi System at 400°C”, Russ. Metall., (5), 6-9 (1982) (Equi.

Diagram, Experimental, #, *, 13)

[1982Hwa] Hwang, C.M., Meichle, M., Salmon, M.B., Wayman, C.M., “Transformation Behavior of

Ti50Ni47Fe3 Alloy: I. Incommensurate and Commensurate Phases”, J. Phys., 43(12),

C4-231-C4-236 (1982) (Crys. Structure, Experimental, 22)

[1982Mat] Matveeva, N.M., Kovneristyi, Yu.K., Savinov, A.S., Sivokha, V.P., Khachin, V.N.,

“Martensitic Transformations in the TiPd-TiNi System”, J. Phys., 43(12), C4-249-C4-253


[1982Mur] Murray, J.L., “The Pd-Ti (Palladium-Titanium) System”, Bull. Alloy Phase Diagram., 3,

321-329 (1982) (Equi. Diagram, Review, #, *, 25)

[1983Siv] Sivokha, V.P., Savinov, A.S., Voronin, V.P., Khachin, V.N., “Martensitic Transformations

and the Shape Memory Effect in Ti0.5Ni0.5-xPdx Alloys”, Phys. Met. Metallogr., 56(3),


[1984Nas] Nash, A., Nash, P., “The Ni-Pd (Nickel-Palladium) System”, Bull. Alloy Phase Diagram, 5,

446-450 (1984) (Equi. Diagram, Review, #, *, 37)

[1985Mat] Matveeva, N.M., Khachin, W.N., Siwokha, W.P., “Diagram of Martensitic Transformations

in System Ni-Pd-Ti” (in Russian), Stable and Metastable Phase Equilibria in Metallic

Systems, Drits, M.E. (Ed.), Nauka, Moscow, 25-29 (1985) (Experimental, Equi. Diagram,

#, *, 3)

[1992Ger] Gerkulova, D.M., Raevskaya, M.V., Tatarkina, A.L., “Isothermal Section of the

Palladium-Titanium-Nickel System at 800°C”, Vestn. Mosk. Univ., Ser. 2: Khim., 33(2),



Martensitic Transformation”, Mater. Trans., JIM, 33(3), 165-177 (1992) (Crys. Structure,

Review, 101)

[1993Mat] Matveeva, N.M., Klopotov, A.A., Kormin, N.M., Sazanov, Yu.A., “Lattice Parameters and

the Sequence of Transformations in Ternary TiNi-TiMe Alloys”, Russ. Metall., (3),

216-220 (1993), translated from Izv. RAN, (3), 233-237 (1993) (Crys. Structure,

Experimental, 10)

[1994Koe] Koetzing, B., “Thermophysical and Thermochemical Properties of Alloys of Ti-Al-Ni,

Ti-Al-Cu, Ti-Al-Pd, Ti-Ni-Pd and Ti-Cu-Pd Ternary Systems” (in German), Final Report

COST 507 I, RWTH Aachen, (1994) (Experimental, 27)

[1997Har] Hara, T., Ohba, T., Okunishi, E., Otsuka K., “Structural Study of R-Phase in Ti-50.23 at.%

Ni and Ti-47.75 at.% Ni-1.50 at.% Fe Alloys”, Mater. Trans., JIM, 38(1), 11-17 (1997)


[1997Sem] Semenova, E.L., “Alloys with the Shape Memory Effect in Systerms of d-Transitioin Metals

Containing Metals of the Platinum Group”, Powder Metall. Met. Ceram., 36(7-8), 394-404

(1997) (Crys. Structure, Review, 36)

[1999Ots] Otsuka, K., Ren, X., “Recent Developments in the Research of Shape Memory Alloys”,

Intermetallics, 7, 511-528 (1999) (Crys. Structure, Review, 131)

[2000Shi] Shiraka, Y., Morizono, Y., Nishida, M., “New Precipitate Phase in Pd and Ni Rich Ti-Pd-Ni

Shape Memory Alloys”, Mater. Sci. Forum, 327-328, 171-174 (2000) (Crys. Structure,

Experimental, 8)

[2000Xu] Xu, H., Hu, S., Gong, S., “Influence of the Recrystallization Process on the Mechanical

Properties of Ti51Ni13Pd36 High Temperature Memory Alloy”, Z. Metallkd., 91(6), 468-471


428


MSIT®

Ni–Pd–Ti

[2001Tia] Tian, Q., Wu, J., “Phase Transformation Behaviour and Microstructure of Ti51Pd30Ni19

Alloy”, Z. Metallkd., 92(5), 436-440 (2001) (Crys. Structure, Experimental, 14)

[2002Tia] Tian, Q., Wu, J., “Tensile Behavior of Ti50.6Pd30Ni19.4 Alloy Under Different Tensile

Conditions”, Mater. Sci. Eng. A, A325, 249-254 (2002) (Experimental, 12)

[2003Tia] Tian, Q., Chen, J., Chen, Y., Wu, J., “Oxidation Behaviour of TiNi-Pd Snape Memory

Alloys”, Z. Metallkd., 94(1), 36-40 (2003) (Crys. Structure, 12)

[2004Ted] Tedenac, J.C., Velikanova, T., Turchanin, M., “Ni-Ti (Nickel-Titanium)”, MSIT Binary

Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), Materials Science

International Services, GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure,

Equi. Diagram, Assessment, 37)


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

(Ni,Pd)

(Ni)

1455

(Pd)

1555

cF4

Fm3m

Cu

a = 352.32

a = 389.01

pure Ni at 20°C [V-C2]

pure Pd at 20°C [V-C2]

( Ti)(h)

1670 - 882

cI2

Im3m

W

a = 330.65 [Mas2]

( Ti)(r)

882

hP2

P63/mmc

Mg

a = 295.06

c = 468.25


Ti4Pd

585

cP8

Pm3m

Cr3Si

a = 500.5 at 20 at.% Pd and 20°C [1982Mur]

[V-C2]

Ti2Pd

960

tI6

I4/mmm

MoSi2

a = 309.0 to 309.5

c = 1005.0 to

1005.4

at 33.3 at.% Pd and 20°C [1982Mur]

dissolves up to 25 mol% Ti2Ni at 800°C

Ti(Ni,Pd)

B2 phase

TiPd

1400

TiNi

1311

cP2

Pm3m

CsCl a = 318.0

a = 299.8

a = 301.0

TiPd at 700°C [1982Mur] homogeneity

range is 47 to 53 at.% Pd [1982Mur]

for Ti56.5Ni43.5

for Ti43Ni57 [2004Ted]

429


MSIT®

Ni–Pd–Ti

Ti(Ni,Pd)

B19 martensite

< 510

oP4

Amma

AuCd

a = 278.0 to 281.0

b = 455.0 to 456.0

c = 486.0 to 489.0

a = 280.0

b = 459.0

c = 484.0

a = 282.0

b = 432.0

c = 460.0

a = 278.0

b = 445.0

c = 451.0

a = 278.0

b = 451.0

c = 477.0

a = 278.0

b = 451.0

c = 477.0

[1982Mur]

Ti50Pd50 at 20°C [1985Mat]

Ti50Ni38.6Pd11.4 at 20°C [1983Siv,

1985Mat]

Ti50Ni24.5Pd25.5 at 20°C [1983Siv]

Ti50Ni13.3Pd36.7 at 20°C [1983Siv]

Ti50Ni5Pd45 at 20°C [1983Siv]

Ti2Pd3

1330

oC20

Amma

Au2V

a = 461.0

a = 1433.0

a = 464.0

at 60 at.% Pd and 20°C [1982Mur]

Ti3Pd5

1290

tP8

P4/mmm

distorted MoSi2

a = 326.3

c = 1143.6

at 62.0 at.% Pd and 20°C [1982Mur]

TiPd2

1400

tI6

I4/mmm

MoSi2

o?

a = 324.0

c = 848.0

a = 341.0

b = 307.0

c = 856.0

at 66.7 at.% Pd and 20°C [1982Mur]

homogeneity range is from 65 to 67 at.%

Pd [1982Mur]

[1982Mur]

Ti(Ni,Pd)3

TiPd3

< 1530

TiNi3 1380

hP16

P63/mmc

TiNi3

a = 548.9 to 548.95

c = 896.4 to 897.39

a = 555.56

c = 933.33

a = 510.28

c = 827.19

a = 522.22

c = 888.89

at 75 at.% and 20°C [1982Mur]

(Ni0.27Pd0.73)3Ti and 20°C [1992Ger]

75 to 80.1 at.% Ni at 1300°C [2004Ted]

[V-C2]

(Ni0.67Pd0.33)3Ti and 20°C [1992Ger]

Ti2Ni

984

cF96

Fd3m

Ti2Ni

a = 1127.8 to

1132.4

33 to 34 at.% Ni [2004Ted]

dissolves up to 12 mol% Ti2Pd at 800°C

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

430


MSIT®

Ni–Pd–Ti

Ti(Ni,Pd) (martensite)

R phase

hP18

P3

-

a = 735.8

c = 528.55

Ni50.23Ti49.77, at 20°C [1997Har].

X-ray diffraction.

Ti(Ni,Pd)

B19´ martensite

mP4

P21/m

TiNi

a = 289.8

b = 410.8

c = 464.6

= 97.78°

a = 288.0

b = 420.0

c = 466.0

= 96.6°

a = 278.0

b = 433.0

c = 464.0

= 90.4°

Ni49.2Ti50.8, at 20°C [1992Shi].

Single crystal X-ray diffraction.

Ti50Ni42.6Pd7.4 at 0°C [1983Siv,

1985Mat]

Ti50Ni36.5Pd13.5 at -130°C [1983Siv]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

10 20 30 40

0

250

500

Ti 50.00

Ni 50.00

Pd 0.00

Ti 50.00

Ni 0.00

Pd 50.00Pd, at.%

Te

mp

era

ture

, °C

B2+B19'

B2

B19

B19'

MS

MF

MS

MF

TR

B2+B19'

R

R+B19'

B2+B19

B2(ω)

Fig. 1: Ni-Pd-Ti.

Structural diagram

for diffusionless

displasive

transformations

along TiNi-TiPd

line. Ms and MF are

martensite start and

martensite finish

temperatures

431


MSIT®

Ni–Pd–Ti

40 30 20 10

1000

1100

1200

1300

1400

1500

Ti 50.00

Ni 50.00

Pd 0.00

Ti 50.00

Ni 0.00

Pd 50.00Ni, at.%

Te

mp

era

ture

, °C

L

L+B2

B2

Fig. 2: Ni-Pd-Ti.

The pseudobinary

section NiTi - PdTi

60

70

80

90

10 20 30 40

10

20

30

40

Ti Ti 50.00

Ni 50.00

Pd 0.00

Ti 50.00

Ni 0.00

Pd 50.00Data / Grid: at.%

Axes: at.%

(βTi)

Ti2Pd

Ti2Ni

B2

(βTi)+Ti2Ni+Ti

2Pd

Ti2Ni+B2+Ti

2Pd

(βTi)+Ti2Pd

(βTi)+Ti2Pd

(βTi)+Ti2Ni

Ti2Ni+B2

(βTi)

(βTi)+B2+Ti2Pd

Fig. 3: Ni-Pd-Ti.

A partial isothermal

section at 900°C

432


MSIT®

Ni–Pd–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni

Pd Data / Grid: at.%

Axes: at.%

Ti(Ni,Pd)3

(Ni,Pd)+TiNi3

TiPd3

TiNi3

(Ni,Pd)

?

60

70

80

90

10 20 30 40

10

20

30

40

Ti Ti 50.00

Ni 50.00

Pd 0.00

Ti 50.00

Ni 0.00


Axes: at.%

Ti2Pd+Ti

2Ni+Ti(Ni,Pd)

Ti2Ni+Ti(Ni,Pd)

(αTi)+Ti2Ni

(βTi)+Ti2Pd

Ti2Pd+Ti(Ni,Pd)

Ti(Ni,Pd)

(βTi)+Ti2Pd+Ti

2Ni

(αTi)+(βTi)Ti

2Ni

Ti2Pd

(αTi)+(βTi)+Ti2Pd

(βTi)

(αTi)

Fig. 4: Ni-Pd-Ti.


section at 800°C

Fig. 5: Ni-Pd-Ti.


section at 700°C

433


MSIT®

Ni–Pd–Ti

10 20 30

300

400

500

600

700

800

900

1000

1100

Ti 66.70

Ni 0.00

Pd 33.30

Ti 66.70

Ni 33.30

Pd 0.00Ni, at.%

Te

mp

era

ture

, °C

L

(βTi)Ti2Pd

1120°C

960°C

Ti2Pd

L+Ti2Pd+Ti2Ni

Ti2NiTi2Pd+Ti2Ni

L+(βTi)+Ti2Pd

L+(βTi)

(βTi) L+TiNi+Ti2Ni

L+TiNi

L+Ti2Ni

1028°C

984°C

60

70

80

90

10 20 30 40

10

20

30

40

Ti Ti 50.00

Ni 50.00

Pd 0.00

Ti 50.00

Ni 0.00


Axes: at.%

(αTi)+Ti2Pd

Ti2Pd+TiNi(?)+Ti

2Ni

Ti2Pd+TiNi(?)

Ti2Pd+Ti

2Ni+(αTi)

Ti2Pd+Ti

2Ni

TiNi(?)

B19+Ti2Pd+TiNi(?)

Ti4Pd

Ti2Ni+TiNi(?)

Ti2Ni

(αTi)+Ti2Ni

Ti2Pd+Ti

4Pd

(αTi)

Ti2Pd

(αTi)+Ti4Pd+Ti

2Pd

Ti2Pd+B19 B19

B19+TiNi(?)

Fig. 7: Ni-Pd-Ti.

An isopleth at a

constant Ti content of

66.67 at.%

Fig. 6: Ni-Pd-Ti.


section at 400°C

434


MSIT®

Ni–Pd–Ti

1200

Ti 0.00

Ni 40.00

Pd 60.00

Ti 10.00

Ni 36.00

Pd 54.00Ti, mass%

Te

mp

era

ture

, °C

L

(Pd)

L+(Pd)

(Pd)+Ti(Ni,Pd)3

L+Ti(Ni,Pd)3

L+(Pd)+Ti(Ni,Pd)3

2 4 6 8

1225

1250

Fig. 8: Ni-Pd-Ti.

An isopleth along the

line of constant Pd:Ni

mass ratio of 3:2

435


MSIT®

Ni–Si–Ti

Nickel – Silicon – Titanium

Nathalie Lebrun

Literature Data

Seven ternary compounds have been reported in the Ni-Si-Ti system. The existence of 5 is not well

established. Some workers attributed this phase to the extension of the martensitic transformation observed

in TiNi alloys into the ternary system.

Isothermal sections at 750, 1000 and 1100°C have been determined by [1958Wes, 1966Mar1, 1971Wil,

1999Hu] using SEM, WDX, X-ray diffraction and electron probe X-ray. Some disagreements are observed

due to the non crystallization of the Ti5Si4 phase. Only [1999Hu] detected this binary phase in a study of

the isothermal section for 1110°C involving 68 alloys (with purity of 99.99 at.% Si, 99.9 at.% Ti and 99.99

at.% Ni). Discrepancies were also noticed concerning the extension of the solubility range of the (Ni) phase

into the ternary system. [1958Wes] reported a large solubility range of up to 20 at.% Ni whereas [1966Mar1,

1971Wil] suggested a smaller solubility range (up to 10 at.% Ni).

[1980Bud] proposed a partial liquidus surface in the Ti rich corner. Thermal analysis, microstructural

observation, microhardness and electrical resistivity measurement were used. A ternary eutectic was found

close to the Ni-Ti binary edge. They also found a pseudobinary eutectic along the Ti5Si3-Ti2Ni section.

Binary Systems

The Ni-Si binary system has been assessed by [1991Nas]. More recently, thermodynamic evaluations were

carried out by [1996Lin, 1999Du, 2001Tok]. Discrepancies are observed between the experimental points

and the calculated phase equilibria, in particular concerning the solid solubility of Si in (Ni). Moreover,

large differences are noted concerning the calculated temperatures of the monovariant reactions.

Consequently, the binary system of [1991Nas] was accepted for this assessment.

The Ni-Ti binary system has been reviewed extensively by [1987Mur1]. More recently, [1996Bel] has

presented a new assessment of the thermodynamic properties of the stable phases based on thermochemical

and phase diagram data available in the literature. Their calculation is in good agreement with the phase

equilibria reported by [1987Mur1]. The solid state homogeneity range of TiNi3 has been reproduced by

[1996Bel] and is in good agreement with the literature data. Using a symmetric two sublattice model,

[1999Tan] described the transformation from the ordered TiNi phase to the disordered TiNi bcc phase,

which occurs at around 93.5°C. The accepted diagram in the present assessment is thus a compilation of

[1987Mur1] and [1996Bel].

A complete evaluation of the Si-Ti system was carried out by [1987Mur2]. A thermodynamic optimization

was carried out by [1996Sei]. Good agreement is observed between the experimental data from the literature

and the calculated phase diagram, except in the Si rich corner where the calculated curves are systematically

lower than the experimental data. Recently, [1998Du] reinvestigated experimentally the Si rich part of the

diagram and determined a eutectic temperature of 1331°C with a liquid composition of 85 at.% Si, in

agreement with the previous experimental data available in the literature. Consequently, the Si-Ti system is

accepted from [1998Du] for the Si-rich part and from [1996Sei] for the composition range 0 to 60 at.% Si.

Solid Phases

Crystallographic data for all of the solid phases in the system are presented in Table 1.

Seven ternary phases have been found in the Ni-Si-Ti system. The 1 phase was first identified by

[1958Wes] at the composition TiNiSi. Later, [1963Spi] determined the crystal structure to be orthorhombic

with a large unit cell (a = 702, b = 518, c = 1111 pm). More extensive work carried out by [1965Sho] on

single crystals resulted in a smaller unit cell. This last result was confirmed by later work [1969Jei,

1999Hu]. The solubility range at 1100°C was reported by [1999Hu]. In the review of [2001Hof], an

436


MSIT®

Ni–Si–Ti

alternative space group to that reported by [1965Sho] was proposed but with the same values for the lattice

parameters. The crystal data of [1965Sho] and [2001Hof] are given in Table 1.

The 2 phase was studied extensively by [1961Bar] using samples that had been annealed for 3 days at

1100°C. This ternary phase has the largest homogeneity range [1999Hu] of all of the ternary compounds in

this system. The phase composition varies from TiNi2Si to Ti2Ni3Si. A large elongation in the lattice

parameters has been observed (1.12 % for a and 3.05 % for c). Later, [2002Hsi2] estimated the solubility

range of 2 to be from 33 to 35 at.% Ti and from 16 to 19 at.% Si at 950°C.

The 3 phase has been reported recently by [1999Hu]. This phase shows little compositional variation.

The 4 phase was first reported by [1958Wes] and was later found to be tetragonal by [1966Mar1]. The

space group was specified by [1966Mar2, 1967Mar, 1969Jei] with 56 atoms in the cell. [1969Jei]

investigated the crystal structure from a sample annealed for 4 days at 1000°C and confirmed the previous

results. From neutron and X-ray diffraction experiments, [1990Hor] confirmed the lattice parameters,

which are reported in Table 1. Nevertheless, disagreement is observed considering the experimental

investigations of [1985Hun] who found an orthorhombic structure with a = 896, b = 805, c = 377 pm with

a composition of TiNiSi2 for 3. The crystal structure was deduced from a sample with a starting

composition of Ti52Ni48 in contact with Si at 600°C. The final composition of Ti25Ni25Si50 was obtained

after annealing for 30 min at 625°C. Only the tetragonal structure has been retained in this assessment since

the sample used by [1985Hun] was certainly not in equilibrium.

The existence of the 5 phase has been disputed. [1958Wes, 1966Mar1, 2002Hsi1, 2002Hsi2] found it

whereas [1999Hu] did not. The composition is found to be close to the homogeneity range of TiNi

[1958Wes]. This was confirmed by [1966Mar1] who deduced a composition range of 50 to 55 at.% Ti and

8 to 12 at.% Si. Later, [2002Hsi2] detected the same phase after homogenization at 950°C for 72 h and

annealing at 900°C for 2 h and confirmed the homogeneity range. They suggested that this ternary

compound was a result of the martensitic transformation in TiNi. [2002Hsi1] found a monoclinic phase and

the lattice parameters are reported on Table 1.

The 6 phase was first reported by [1958Wes]. Its existence was confirmed later [1956Bea, 1957Bea,

1963Spi, 1966Mar2, 1971Wil, 1999Hu]. 6 is found to have a quite large homogeneity range at 1100°C

[1999Hu], which becomes smaller at lower temperatures [1971Wil]. The existence of a measurable

homogeneity range explains the slight discrepancies in the composition given in the literature: Ti2Ni7Si3[1958Wes], TiNi2Si [1956Bea, 1957Bea], Ti6Ni16Si7 [1963Spi, 1999Hu], Ti20.7Ni58.3Si23 [1971Wil].

The 7 phase has been systematically detected [1958Wes, 1966Mar1, 1974Ste, 1999Hu]. [1974Ste]

investigated its crystal structure. Data are reported in Table 1. The space group has not been well defined.

[1974Ste] suggested several possibilities for the space group: P622, P6mm, P62m, P6m2 or P6/mmm.


[1980Sha] reported investigations carried out by [1980Bud] in the Ti-rich part of the ternary system. The

section Ti5Si3-Ti2Ni was found to be pseudobinary with a L Ti5Si3 + Ti2Ni eutectic reaction occurring at

960°C and a composition of 63Ti-34Ni-3Si (mass%) [1980Bud].


A partial reaction scheme is reported in Fig. 1. [1980Bud] observed a ternary eutectic E. The compositions

of the different phases in the eutectic are reported in Table 2. In addition, the liquid undergoes a eutectic

decomposition L Ti5Si3 + Ti2Ni at 960°C (reaction e2 in Fig. 1). The liquid composition is reported in

Table 2. From experimental investigations, [1980Bud] proposed solid state reactions in the ternary system,

which have also been indicated in the reaction scheme (Fig. 1).

Liquidus Surface

The partial liquidus surface in the Ti rich part of the system is taken from [1980Bud] and reported in Fig. 2.

It was first estimated by mathematical calculation and was then confirmed by thermal analysis and

microstructural observation of annealed samples.

437


MSIT®

Ni–Si–Ti

Isothermal Sections

[1980Bud] reported a maximum solubility of Si and Ni in ( Ti) into the ternary system of 0.3 and

0.2 mass%, respectively. The results are in agreement with [1966Mar1].

Disagreement is observed in the Ni rich part of the isothermal section for 1000°C as determined by

[1966Mar1] and [1971Wil]. [1971Wil] proposed a less extensive homogeneity range for the (Ni), and the

1 phase extends to higher Ti content than given in [1958Wes]. Later, [1999Hu] confirmed the homogeneity

range of 1 found by [1971Wil]. The homogeneity range of (Ni) measured by [1971Wil] is in agreement

with that indicated by [1966Mar1] for 750°C. Consequently, the partial isothermal section presented by

[1971Wil] for 1000°C has been retained in this assessment.

Discrepancies are also noted concerning the phase equilibria involving TiSi, Ti5Si4, Ti5Si3 and the ternary

compounds 1 and 4. Only [1999Hu] observed the crystallization of Ti5Si4, which exists below 1920°C.

Thermal equilibrium may not have been reached in the samples investigated by [1958Wes, 1966Mar1,

1992Lut] because of the slow kinetics of crystallization of Ti5Si4 at low temperatures.

Moreover, the Ti3Si phase is not indicated in all the isothermal sections presented while this phase exists

down to 1170°C. [1999Hu] suggests that Ti3Si may have decomposed during heat treatment at 500°C by

slight oxidation.

Consequently, the missing phases have been added to the isothermal sections presented in Figs. 3 to 5. The

undetermined and deduced phase fields are indicated as dashed lines in the figures. The boundaries of the

binary phases are positioned according the accepted binary phase diagrams. Moreover, the binary phases

NiTi2, Ti3Si, Ti5Si4, TiSi, TiSi2, , , NiSi, and ´ were considered as stoichiometric phases in agreement

with the binaries since the homogeneity ranges of these phases reported by [1958Wes, 1966Mar1, 1971Wil,

1999Hu] are only speculative.


Microhardness measurements have been reported by [1971Wil]. The characteristic microhardness values

for the phases after heat treatment were approximately 350 HV for (Ni), 425 HV for Ni3Si, 500 HV for

Ni3Ti, 820 HV for 6, 900 HV for Ni5Si2 and 1000 HV for Ni2Si. [1969Wil] also reported tensile data for

Ni-Si-Ti alloys in the Ni rich corner with 3.1 to 13.3 mass% Si and 0.85 to 4.2 mass% Ti. A constant 0.1%

proof stress was observed at about 417 MPa, whereas an increase of the UTS from 618 to 927 MPa at 900°C

and from 896 to 973 MPa was noticed during the annealing time. [2002Hsi1] studied the effects of aging,

cold rolling and thermal cycling on the martensitic transformation of a Ti51Ni47Si2 alloy.

Miscellaneous

[1989Set] suggested that at 550°C, Si becomes mobile and diffuses into the Ni-Ti compound, resulting in

the growth of the ternary phase 4. If it is in excess with respect to this ternary silicide, a separate layer of

Ni silicide grows between the substrate and 4, owing to the fact that Ni is the main diffusing species.

[1985Hun] found that the TiNi compound, when in contact with Si, appeared to be relatively stable under

annealing. The reaction between the TiNi phase and the Si substrate takes place at 600°C via the migration

of Si into the alloy matrix. Annealing at 625°C for 30 min results in a reacted layer of Ti25Ni25Si50. It was

found that a 7 % Si addition to TiNi composites increases the melting temperature by 50°C [1984Che].

[1955Rob] indicated that the silicide Ti5Si3 did not react with molten nickel.

References

[1955Rob] Robins, D.A., Jenkins, I., “The Stability of Some metal Silicides of Potential Value in High

Temperature Materials”, Plansee Proc., Reutte, Tyrol, 187-198 (1955) (Experimental, 9)

[1956Bea] Beattie, H.J., VerSnyder, F.L., “A New Complex Phase in a High-Temperature Alloy”,

Nature, 178, 208-209 (1956) (Experimental, Crys. Structure, 0)

[1957Bea] Beattie, H.J., Hagel, W.C., “Intermetallic Compounds in Titanium-Hardened Alloys”,

J. Met. Trans. AIME, 209, 911-917 (1957) (Experimental, Crys. Structure, Mechan.

Prop., 8)

438


MSIT®

Ni–Si–Ti

[1958Wes] Westbrook, J.H., Di Cerbo, R.K., Peat, A.J., Rep. 58-rc-2117, General Electronic Research

Laboratory (1958) cited in [1999Hu]

[1961Bar] Bardos, D.I., Gupta, K.P., Beck, P.A., “Ternary Laves Phases with Transition Elements and

Silicon”, Trans. Metall. Soc. AIME, 221, 1087-1088 (1961) (Experimental, Crys.

Structure, 10)

[1961Gla] Gladyshevsky, E.I., Kripyakevich, P.I., Kuzma, Y.B., Teslyuk, M.Y., “New

Representatives of Structures of Mg6Cu16Si7 and Th6Mn23”, Kristallografiya, 6, 769-770


[1962Gla] Gladyshevskii, E.I., Franko, I.Y., “Crystal Structure of Compounds and Phase Equilibria in

Ternary Systems of Two Transition Metals and Silicon”, Inorg. Mater., 4, 262-265 (1962),

translated from Poroshk. Metall., 4(10), 46-49 (1962) (Experimental, Crys. Structure, 17)

[1963Spi] Spiegel, F.X., Bardos, D., Beck, P.A., “Ternary G and E Silicides and Germanides of

Transition Elements”, Trans. Metall. Soc. AIME, 227, 575-579 (1963) (Experimental, Crys.

Structure, 13)

[1965Sho] Shoemaker, C.B., Shoemaker, D.P., “A Ternary Alloy with PbCl2-Type Structure:

TiNiSi(E)”, Acta Crystallogr., 18, 900-905 (1965) (Crys. Structure, Experimental, 20)

[1966Mar1] Markiv, V.Y., Gladyshevskii, E.I., Kripyakevich, P.I., Fedoruk, T.I., “The System

Titanium-Nickel-Silicon”, Inorg. Mater., 2, 1126-1128 (1966) translated from Izv. Akad.

Nauk SSSR, Neorg. Mater., 2, 1317-1319 (1966) (Experimental, Equi. Diagram, Crys.

Structure, 23)

[1966Mar2] Markiv, V.J., “The Crystal Structures of the Compounds R(M, X)2 and RMX2 in Zr-Ni-Al,

Ti-Fe-Si and Related Systems”, Acta Crystallogr., 21, A84 (1966) (Abstract)

[1967Mar] Markiv, V.Y., Gladyshevsky, E.I., Skolozdra, R.V., Kripyaakevich, “Ternary Compounds

RX´X” in the Systems Ti-V(Fe, Co, Ni)-Si and Some Related Systems”, Dop. Akad. Nauk

Ukr. RSR, A3, 268 (1967) (Abstract)

[1969Jei] Jeitschko, W., Jordan, A.G., Beck, P.A., “V and E Phases in Ternary Systems with

Transition Metals and Silicon or Germanium”, Trans. Metall. Soc. AIME, 245, 335-339

(1969) (Experimental, Crys. Struct., 27)

[1969Wil] Williams, K.J., “The Microstructure and Tensile Properties of Nickel-Rich Nickel-Silicon

and Nickel-Silicon-Titanium Alloys”, J. Inst. Met., 97, 112-118 (1969) (Experimental,

Equi. Diagram, Mechan. Prop., 21)

[1971Wil] Williams, K.J., “The 1000°C (1273 K) Isotherm of the Ni-Si-Ti System from 0 to 16 % Si

and 0 to 16 % Ti”, J. Inst. Met., 99, 310-315 (1971) (Experimental, Equi. Diagram, 8)

[1974Ste] Steinmetz, J., Albrecht, J.M., Malaman, B., “A New Family of Ternary Silicides of the

General Formula TT´4Si3 (T = Ti, Nb, Ta; T´ = Fe, Co, Ni)”, Compt. Rend. Acad. Sci. Paris,

279C, 1119-1120 (1974) (Experimental, Crys. Structure, 4)

[1980Bud] Budberg, P.B., Alisova, S.P., Kobilkin, A.N., “Phase Equilibria in the Ternary System Ti-

Ti5Si3-Ti2Ni”, Dokl. Akad. Nauk USSR, 250(5), 1137-1140 (1980) (Experimental, Equi.

Diagram, 7)

[1980Sha] Shaling, R.E., Kovneristy, Y.K., “Phase Stability and Phase Equilibrium in Titanium

Alloys”, Metall. Soc. AIME. 420 Commonwealth Dr., Warrendale Pens. 15086, 72-306,


[1984Che] Chepeleva, V.P., Delevi, V.G., Trunevich, L.V., Manzheleeva, T.N., Demyanchuk, A.F.,

“Dopping of Titanium-Nickel Powder Compositions of the Eutectic Composition”,

Poroshk. Metall., 2, 57-60 (1984) (Experimental, 4)

[1985Hun] Hung, L.S., Mayer, J.W., “Interactions of Four Metallic Compounds with Si Substrates”,

J. Appl. Phys., 60(3), 1002-1008 (1986) (Experimental, Crys. Structure, 21)

[1987Mur1] Murray, J.L., “The Ni-Ti (Nickel-Titanium) System” in “Phase Diagrams of Binary



439


MSIT®

Ni–Si–Ti

[1987Mur2] Murray, J.L., “The Si-Ti (Silicon-Titanium) System” in “Phase Diagrams of Binary



[1989Set] Setton, M., Spiegel, J., Rothman, B., “Formation of a Ternary Silicide for Ni/Ti/Si (100) and

Ni/TiSi2 Structures”, J. Mater. Res., 4(5), 1218-1226 (1989) (Experimental, 16)

[1990Hor] Horache, E., Feist, T.P., Stuart, J.A., Fischer, J.E., “Neutron Rietveld Refinement of the

Structure of the Ternary Silicides Ti4Ni4Si7”, J. Mater. Res., 5(9), 1887-1893 (1990)


[1991Nas] Nash, P., Nash, A., “Ni-Si (Nickel-Silicon)” in “Phase Diagrams of Binary Nickel Alloys”,

ASM Internal, Metals Park, OH, 299-306 (1991) (Equi. Diagram, Crys. Structure,


[1992Lut] Lutskaya, N.V., Alisova, S.P., “Phase Structure of TiCu(TiNi, TiCo)-TiSi Sections in

Ternary Ti-Cu(Ni, Co)-Si Systems”, Russ. Metall., 3, 180-182 (1992) translated from Izv.

Akad. Nauk SSSR Met., 3, 194-196 (1992) (Equi. Diagram, Experimental, 9)

[1996Bel] Bellen, P., Kumar, K.C.H., Wollants, P., “Thermodynamic Assessment of the Ni-Ti Phase

Diagram”, Z. Metallkd., 87(12), 972-978 (1996) (Assessment, Equi. Diagram,

Thermodyn., 43)

[1996Lin] Lindholm, M., Sundman, B., “A Thermodynamic Evaluation of the Nickel-Silicon System”,

Metall. Mater. Trans. A, 27A, 2897-2903 (1996) (Equi. Diagram, Thermodyn., 23)

[1996Sei] Seifert, H.J., Lukas, H.L., Petzow, G., “Thermodynamic Optimization of the Si-Ti System”,

Z. Metallkd., 87, 2-13 (1996) (Equi. Diagram, Thermodyn., 63)

[1998Du] Du, Y., Schuster, J.C., “A Reinvestigation of the Constitution of the Partial System SiTi-

Si”, J. Mater. Sci. Lett., 17, 1407-1408 (1998) (Equi. Diagram, Experimental, 7)

[1999Du] Du, Y., Schuster, C., “Experimental Investigations and Thermodynamic Descriptions of the

Ni-Si and C-Ni-Si Systems”, Metall. Mater. Trans. A, 30A, 2409-2418 (1999) (Equi.

Diagram, Experimental, Thermodyn., 44)

[1999Hu] Hu, X., Chen, G., Ion, C., Ni, K., “The 1100°C Isothermal Section of the Ti-Ni-Si Ternary

System”, J. Phase Equilib., 20(5), 508-514 (1999) (Experimental, Equi. Diagram, Crys.

Structure, 25)

[1999Tan] Tang, W., Sundman, B., Sandstroem, R., Qiu, C., “New Modelling of the B2 Phase and its

Associated Martensitic Transformation in the Ni-Ti System”, Acta Mater., 47(12), 3457-

3468 (1999) (Thermodyn., Calculation, 51)

[2001Hof] Hoffmann, R.D., Pöttgen, R., “AlB2-Related Intermetallic Compounds - a Comprehensive

View Based on Group-Subgroup Relations”, Z. Kristallogr., 216, 127-145 (2001) (Review,


[2001Tok] Tokunaga, T., Nishio, K., Hasebe, M., “Thermodynamic Study of Phase Equilibria in the

Ni-Si-B System”, J. Phase Equilib., 22(3), 291-299 (2001) (Equi. Diagram, Thermodyn.26)

[2002Hsi1] Hsieh, S.F., Wu, S.K., Lin, H.C., Martensitic Transformation of a Ti-rich Ti51Ni47Si2 Shape

Memory Alloy”, J. Alloys Compd., 335, 254-261 (2002) (Experimental, Crys. Structure, 27)

[2002Hsi2] Hsieh, S.F., Wu, S.K., Lin, H.C., “Transformation Temperatures and Second Phases in Ti-

Ni-Si Ternary Shape Memory Alloys with Si 2 at.%”, J. Alloys Compd., 339, 162-166


440


MSIT®

Ni–Si–Ti


Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

( Ti)

1670 - 882

cI2

Im3m

W

a = 330.65 pure Ti at 25°C [Mas2]

dissolves up to 10.6 at.% Ni at 942°C

[1996Bel]

dissolves up to 5 at.% Si at 1340°C

[1996Sei]

( Ti)

< 882

hP2

P63/mmc

Mg

a = 295.06

c = 468.35


dissolves up to 1 at.% Ni at 767°C

[1996Bel]

dissolves up to 1.11 at.% Si at 862°C

[1996Sei]

(Ni)

< 1455

cF4

Fm3m

Cu

a = 352.40 pure Ni at 25°C [Mas2]

dissolves up to 15.30 at.% Ti at 1300°C

[1996Bel]

Dissolves up to 15.8 at.% Si at 1143°C

[1991Nas]

(Si)

< 1414

cF8

Fd3m

C (diamond)

a = 543.06 pure Si at 25°C [Mas2]

100 at.% Si [1991Nas]

100 at.% Si [1996Sei]

TiNi3< 1380

hP16

P63/mmc

TiNi3

a = 510.28

c = 827.19

75 to 80.1 at.% Ni at 1300°C [1996Bel]

[V-C2]

TiNi

< 1310

cP2

Pm3m

CsCl

a = 301.0

49.5 to 57 at.% Ni [1987Mur1]

Ti0.98Ni1.02 [V-C2]

Ti2Ni

< 984

cF96

Fd3m

Ti2Ni

a = 1132.4

33 to 34 at.% Ni [1987Mur1]

annealed at 950°C for 72 hours [V-C2]

Ti5Si3< 2130

hP16

P63/mcm

Mn5Si3

a = 746.10 0.03

c = 515.08 0.01

35.9 to 39.3 at 1340°C [1996Sei]

[V-C2]

Ti5Si4< 1920

tP36

P41212

Zr5Si4

a = 713.3

c = 1297.7

at 44.44 at.% Si [1996Sei]

[1987Mur2]

TiSi

< 1570

oP8

Pnma or Pmm2

FeB or TiSi

a = 655.1 0.6

b = 363.3 0.3

c = 498.3 0.5

or

a = 361.8

b = 649.2

c = 497.3

at 50 at.% Si [1996Sei]

two possible space group [1987Mur2]

TiSi2< 1488

oF24

Fddd

TiSi2

a = 826.71 0.09

b = 480.00 0.05

c = 855.05 0.11

at 33.33 at.% Si [1996Sei]

[1987Mur2]

441


MSIT®

Ni–Si–Ti

Ti3Si

< 1170

tP32

P42/n

Ti3P

a = 1020.6 0.6

c = 506.9 0.2

at 25 at.% Si [1996Sei]

[1987Mur2]

, Ni2Si

1306-825

hP6

P6322

Ni2Si

a = 383.6 0.1

c = 494.8 0.1

33.4 to 41 at.% Si

Ni1.86Si1.14 [V-C2]

, Ni2Si

< 1255

oP12

Pnma

Co2Si

a = 502.2 0.1

b = 374.1 0.1

c = 708.8 0.1

at 33.33 at.% Si

annealed at 650°C for 12 h [V-C2]

, Ni31Si12

< 1242

hP43

P321

Ni31Si12

a = 0.667 0.1

c = 1.228 0.8

26.5 to 29.5 at.% Si [1991Nas]

3, Ni3Si

1170 - 1115

mC16

a = 704 7

b = 626 4

c = 508 1

= 48.84°

24.5 to 25.5 at.% Si

[1991Nas]

2, Ni3Si

1115 - 990

mC16

a = 697 2

b = 625 4

c = 507 8

= 48.74°

24.5 to 25.5 at.% Si

[1991Nas]

1, Ni3Si

< 1035

cP4

Pm3m

AuCu3

a = 351.0

22.8 to 24.5 at.% Si

annealed at 1000°C for 5 days [V-C2]

NiSi

< 992

oP8

Pnma

MnP

a = 562.8 0.2

b = 519.0 0.1

c = 333.0 0.1

at 50 at.% Si

annealed at 800°C for 17 hours [V-C2]

´, NiSi2993 - 981

- - at 66.67 at.% Si [1991Nas]

, NiSi2< 981

cF12

Fm3m

CaF2

a = 540.6 at 66.67 at.% Si [1991Nas]

´, Ni3Si2845 - 800

- - 39 to 41 at.% Si [1991Nas]

, Ni3Si2< 830

oP80 - 39 to 41 at.% Si [1991Nas]

* 1, TiNiSi oP12

Pnma or Pmcn

Co2Si or PbCl2

a = 614.84 0.12

b = 701.73 0.14

c = 366.98 0.07

a = 366.98

b = 701.73

c = 614.84

called E phase

space group Pcma [1965Sho]

space group Pmcn [2001Hof]

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

442


MSIT®

Ni–Si–Ti


* 2, Ti2Ni3Si hP12

P63/mmc

MgZn2

a = 479

c = 755

called G´ phase except in [1966Mar1]

( 1)

33-35 at.% Ti and 16-19 at.% Si at

950°C [2002Hsi2]

at 1000°C [1961Bar]

* 3, Ti4NiSi4 called phase H in [1999Hu]

small homogeneity range

* 4, Ti4Ni4Si7 tI56

I4/mmm

Co4Ge7Zr4 a = 1252.29

c = 493.43

called V phase except in [1966Mar1,

1992Lut] ( 2)

[1990Hor]

* 5, Ti6Ni5Si

a = 284

b = 412

c = 418

= 98°

50-55 at.% Ti and 12-15 at.% Si

[2002Hsi2]

called F phase except in [1966Mar1,

1966Mar2, 2002Hsi1, 2002Hsi2] (X)

Ti53Ni37Si10 [1966Mar1, 1966Mar2]

Ti2Ni3Si [2002Hsi1, 2002Hsi2]

[2002Hsi1]

* 6, Ti6Ni16Si7 cF116

Fm3m

Mn23Th6

a = 1122 1

called G phase except in [1966Mar1] (T)

Ti2Ni7Si3 [1958Wes]

Mg6Cu16Si7 [1966Mar2, 1962Gla]

[1961Gla, 1962Gla]

* 7, Ti14Ni49Si37 hP168

a = 1717.3

c = 786.1

called G” phase except in [1966Mar1]

( 3)

[1974Ste]


Ni Si Ti

L ( Ti) + Ti5Ni3 + Ti2Ni 920 E L

( Ti)

Ti5Ni3Ti2Ni

19.24

0.00

37.50

33.33

4.38

0.00

62.50

0.00

76.38

100

0.00

66.67

L Ti5Ni3 + Ti2Ni 960 e2 (max) L

Ti5Ni3Ti2Ni

28.64

37.50

33.33

5.32

62.50

0.00

66.04

0.00

66.67

Phase/

Temperature Range

[°C]

Pearson Symbol/

Space Group/

Prototype

Lattice Parameters

[pm]

Comments/References

443


MSIT®

Ni–Si–Ti

Fig. 1: Ni-Si-Ti. Partial reaction scheme

Si-Ti A-B-CNi-Si-Ti Ni-Ti

l (βTi)+Ti5Si

3

1340 e1

L Ti5Si

3+Ti

2Ni

960 e2, max

l (βTi)+Ti5Si

3+Ti

2Ni920 E

(βTi)+Ti5Si

3+Ti

2Ni

l + TiNi Ti2Ni

985 p2

(βTi)+Ti5Si

3Ti

3Si

1170 p1

l (βTi) + Ti2Ni

942 e3

(βTi)+Ti5Si

3Ti

3Si+Ti

2Ni890 U

1

Ti5Si

3+Ti

2Ni+Ti

3Si

(βTi)+Ti3Si+Ti

2Ni(βTi) (αTi)+Ti

3Si

862 e4

(βTi)+Ti3Si (αTi)+Ti

2Ni780 U

2

(αTi)+Ti3Si+Ti

2Ni

(αTi)+(βTi)+Ti2Ni

(βTi)+Ti5Si

3+Ti

3Si

(βTi)+Ti3Si+(αTi)

60

70

80

90

10 20 30 40

10

20

30

40

Ti Ti 50.00

Ni 50.00

Si 0.00

Ti 50.00

Ni 0.00


Axes: at.%

e1

e3

E e2

Ti2Ni

Ti5Si

3

Fig. 2: Ni-Si-Ti.

Partial liquidus

surface

444


MSIT®

Ni–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni


Axes: at.%

(Ni)

(Si)

(αTi)

TiNi3

TiNiTi2Ni

Ti3Si

Ti5Si

3

Ti5Si

4

TiSi

TiSi2

ζ

NiSi

εδ

γβ

1

τ4

τ7

τ6

τ2

τ1

τ3

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni


Axes: at.%

(Ni)

(Si)

(βTi)

Ti3Si

Ti5Si

3

Ti5Si

4

TiSi

TiSi2

L

θδ

γ β2

β1

L

TiNi TiNi3

τ4

τ3

τ1

τ7

τ6

τ2

Fig. 3: Ni-Si-Ti.


750°C

Fig. 4: Ni-Si-Ti.


1000°C

445


MSIT®

Ni–Si–Ti

20

40

60

80

20 40 60 80

20

40

60

80

Ti Ni


Axes: at.%

(Ni)

(Si)

(βTi)

Ti3Si

Ti5Si

3

Ti5Si

4

TiSi

TiSi2

L

θδ

γβ

2

TiNi3

TiNiL

τ1

τ2

τ3

τ4

τ6

τ7

Fig. 5: Ni-Si-Ti.


1100°C

light metal systems. part 4: selected systems from al-si-ti to ni-si-ti

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