light metal systems. part 4: selected systems from al-si-ti to ni-si-ti
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XI
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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.
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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
Liquidus, Solidus, Solvus Surfaces
Invariant Equilibria
Pseudobinary Systems
Solid Phases
Binary Systems
Text
References
Tables anddiagrams
Temperature-Composition Sections
Temperature-Composition Sections
Thermodynamics
Notes on Materials Properties and Applications
Thermodynamics
Notes on Materials Properties and Applications
Fig. 1: Structure of a system report
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Introduction
Pseudobinary Systems
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”.
Invariant Equilibria
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.
Liquidus, Solidus, Solvus Surfaces
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.
Notes on Materials Properties and Applications
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.
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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
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rst
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ute
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rnar
y m
axim
um
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tion
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atu
re
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26
1°C
mo
no
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iant
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ibri
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ble
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w
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s
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t 144°C
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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
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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
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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)
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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
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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.
Invariant Equilibria
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.
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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
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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.
Temperature – Composition Sections
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.
Notes on Materials Properties and Applications
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
Landolt-BörnsteinNew Series IV/11A4
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)
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AIME, 212, 60-63 (1958) (Equi. Diagram, Experimental, #, 8)
[1961Bru] Brukl, C., Nowotny, H., Schob, O., Benesovsky, F., “The Crystal Structure of TiSi,
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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.,
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Naturwissenschaften, 50, 41 (1963) (Crys. Structure, Review, 0)
[1965Ram] Raman, A., Schubert, K., “On the Constitution of Some Alloy Series Related to TiAl3. II.
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(1965) (Crys. Structure, Equi. Diagram, #, *, 13)
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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)
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and Ti with Al” (in Russian), Vestn. Mosk. Univ., Ser. Khim., 25, 500-503 (1984) (Equi.
Diagram, Experimental, 13)
6
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Diagram, Experimental, 2)
[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.,
Experimental, #, 48)
[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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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, pseudo- tetragonal
[1961Bru] at Ti31Al19Si50 (x 0.28)
[1965Ram]
Phase/
Temperature Range
[°C]
Pearson Symbol/
Space Group/
Prototype
Lattice Parameters
[pm]
Comments/References
11
Landolt-BörnsteinNew Series IV/11A4
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
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Landolt-BörnsteinNew Series IV/11A4
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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
?
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Landolt-BörnsteinNew Series IV/11A4
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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.%
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
Si 70.00Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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.
Isothermal section at
1200°C
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
Ti3Si
(αTi) α2 TiAl TiAl
2 TiAl3(r)
L
τ1
τ2
possible range of τ3
(Si)Fig. 7: Al-Si-Ti.
Isothermal section at
700°C
15
Landolt-BörnsteinNew Series IV/11A4
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
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Landolt-BörnsteinNew Series IV/11A4
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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].
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Landolt-BörnsteinNew Series IV/11A4
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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.
Temperature – Composition Sections
[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
Landolt-BörnsteinNew Series IV/11A4
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].
Notes on Materials Properties and Applications
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
Landolt-BörnsteinNew Series IV/11A4
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)
[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. 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
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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.
Structure, Experimental, 41)
[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)
[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) (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)
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Sn 60.00Data / Grid: at.%
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
Sn 80.00Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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
Sn 60.00Data / Grid: at.%
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
Sn 60.00Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Invariant Equilibria
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.
Liquidus and Solidus Surfaces
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
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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
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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].
Temperature – Composition Sections
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].
Notes on Materials Properties and Applications
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
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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.
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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
Landolt-BörnsteinNew Series IV/11A4
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Al–Ti–V
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Landolt-BörnsteinNew Series IV/11A4
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Al–Ti–V
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Type Ti-Al-V Compound”, Scr. Metall. Mater., 24, 485-490 (Experimental, 10)
[1991Moh] Mohandas, E., Bearen, P.A., “Site Occupation of Nb, V, Mn and Cr in -TiAl”, Scr. Metall.
Mater., 25, 2023-2027 (1991) (Crys. Structure, Experimental, 15)
[1991Pak] Pak, J.S.L., Lei, C.Y., Wayman, C.M., “Atomic Ordering in Ti-V-Al Shape Memory
Alloys”, Mater. Sci. Eng., A132, 237-244, (1991) (Crys. Structure, Experimental, 6)
[1991Par] Paruchuri, M., Massalski, T.B., “Phase Diagram Relationships in the Ternary System
Ti-Al-V” (in English), Mat. Res. Soc. Symp. Proc., 213, 143-149 (1991) (Equi. Diagram,
Experimental, 19)
[1991Spa] Sparks, C.J., Porter, W.D., Schneibel, J.H., Oliver, W.C., Golec, C.G., “Formation of Cubic
L12 Phases from Aluminum Titanium (Al3Ti) and Aluminum Zirconium (Al3Ar) by
Transition Metal Substitutions for Aluminum”, Mater. Res. Soc. Symp. Proc., 186, 175-80
(1991) (Crys. Structure, Experimental, 15)
[1992Ahm] Ahmed, T., Flower, H.M., “The Phase Transformation in Alloys Based on Titanium
Aluminides Ti3Al-V and TiAl-V”, Mater. Sci. Eng., A152, 31-36 (1992) (Equi. Diagram,
Experimental, #, *, 14)
37
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Al–Ti–V
[1992Cha1] Chaudhury, P.K., Rack, H.J., “Ti-Al-V Ternary Phase Stability at Elevated Temperatures”,
Scr. Metall. Mater., 26, 691-695 (1992) (Crys. Structure, Equi. Diagram, Experimental, 16)
[1992Cha2] Chaudhury, P.K., Long, M., Rack, H.J., “Effect of Vanadium on Elevated Temperature
Phase Relations in Titanium Aluminides Containing 44 at.% Al”, Mater. Sci. Eng., A152,
37-40 (1992) (Equi. Diagram, Experimental, 11)
[1992Gru1] Grujicic, M., Narayan, C.P., “A Study of ´´ Martensitic Transformation Volume
Change in Ti-Al-V Alloys”, Mater. Sci. Eng., 151A, 217-226 (1992) (Crys. Structure,
Experimental, 10)
[1992Gru2] Grujicic, M., “Design of Ti-Al-V Phase for Transformation Toughening of -Titanium
Aluminide”, Mater, Sci. Eng., 154A, 75-78 (1992) (Theory, 12)
[1992Gru3] Grujicic, M., Narayan, C.P., “Effect of Iron Additions on -Phase Stability with Respect to
´´-Martensite in V-28Al-17Ti (mass%)”, Mater. Sci. Eng., 151A, 227-233 (1992)
(Experimental, 9)
[1992Kat] Kattner, U.R., Lin, J.-C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the
Ti-Al System”, Metall. Trans., 23A, 2081-2090 (1992) (Equi. Diagram, Review, Theory,
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[1992Kim] Kimura, M., Hashimoto, K., Morikawa, H., “Study on Phase Stability in Ti-Al-X Systems
at High Temperatures”, Mater. Sci. Eng., 152A, 54-59 (1992) (Equi. Diagram,
Experimental, 12)
[1992Nak] Naka, S., Thomas, M., Khan, T., “Potential and Prospects of Some Intermetallic Compound
for Structural Applications”, Mater. Sci. Technol., 8, 291-298 (1992) (Review, 26)
[1992Naz] Nazmy, M., Staubli, M., Anton, D., “Aspects on Mechanical Behavior of the Ti-Aluminide
Base Intermetallics”, Scr. Metall. Mater., 26, 105-108 (1992) (Experimental, 15)
[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.
Diagram, Experimental, 18)
[1993Hay] Hayes, F.H., “Aluminium - Titanium - Vanadium”, MSIT Ternary Evaluation Program, in
MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services
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,
737-740 (1993) (Crys. Structure, Equi. Diagram, Experimental, 16)
[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)
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[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
Landolt-BörnsteinNew Series IV/11A4
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.
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[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.
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,
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,
187-199 (1997) (Equi. Diagram, Experimental, 17)
[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
Landolt-BörnsteinNew Series IV/11A4
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)
(Crys. Structure, Equi. Diagram, Experimental, 34)
[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)
[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, 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)
[2001Bra] Braun, J., Ellner, M., “Phase Equilibria Investigations on the Aluminium-Rich Part of the
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,
257-261 (2001) (Equi. Diagram, Experimental, 9)
[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
Landolt-BörnsteinNew Series IV/11A4
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,
Equi. Diagram, Experimental, 37)
[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)
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
summarizes several phases [2003Sch1]:
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Al Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Al–Ti–V
20
40
60
80
20 40 60 80
20
40
60
80
Ti V
Al Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Al–Ti–V
20
40
60
80
20 40 60 80
20
40
60
80
Ti V
Al Data / Grid: at.%
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
Al Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Al–Ti–V
20
40
60
80
20 40 60 80
20
40
60
80
Ti V
Al Data / Grid: at.%
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
Al Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Al–Ti–V
20
40
60
80
20 40 60 80
20
40
60
80
Ti V
Al Data / Grid: at.%
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
Al Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Al–Ti–V
20
40
60
80
20 40 60 80
20
40
60
80
Ti V
Al Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Al 25.00Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Temperature – Composition Sections
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
Landolt-BörnsteinNew Series IV/11A4
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.
Notes on Materials Properties and Applications
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
Landolt-BörnsteinNew Series IV/11A4
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,
171-176 (1968) (Equi. Diagram, Experimental, 4)
[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
(1969) (Equi. Diagram, Experimental, 15)
58
Landolt-BörnsteinNew Series IV/11A4
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.
Structure, Experimental, 10)
[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)
[1992Kat] Kattner, U.R., Lin, J.-C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the
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,
737-740 (1993) (Crys. Structure, Experimental, 16)
[1993Oka1] Okamoto, H., “Al-Zr (Aluminium - Zirconium)”, J. Phase Equilib., 14, 259-260 (1993)
(Equi. Diagram, Review, #, 2)
59
Landolt-BörnsteinNew Series IV/11A4
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,
Experimental, #, 15)
[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)
[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, 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)
[2001Bra] Braun, J., Ellner, M., “Phase Equilibria Investigations on the Aluminium-Rich Part of the
Binary System Ti-Al”, Metall. Mater. Trans., 32A, 1037-1047 (2001) (Crys. Structure,
Equi. Diagram, Experimental, 34)
[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,
257-261 (2001) (Crys. Structure, Experimental, 9)
[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
MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services
GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram,
Assessment, 151)
60
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Al–Ti–Zr
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
summarizes several phases [2003Sch2]:
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]
in the (Ti0.84Zr0.16)Al3 alloy at 600°C
[2003Kar]
in the (Ti0.84Zr0.16)Al3 alloy at 700°C
[2003Kar]
Phase/
Temperature Range
[°C]
Pearson Symbol/
Space Group/
Prototype
Lattice Parameters
[pm]
Comments/References
63
Landolt-BörnsteinNew Series IV/11A4
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]
at x = 0.5, as melt-spun ribbon of
2.1Zr-1.1Ti-Al (mass%) [2000Mal]
at x = 0.75, as melt-spun ribbon of
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
Landolt-BörnsteinNew Series IV/11A4
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]
in the as cast (Ti0.68Zr0.32)Al3 alloy
(D023+D022) [2003Kar]
in the as cast (Ti0.76Zr0.24)Al3 alloy
(D023+D022) [2003Kar]
in the mechanically alloyed
Al-8Ti-3.8Zr (mass%) ((Ti0.8Zr0.2)Al3)
[1993Lee]
in the (Ti0.84Zr0.16)Al3 alloy at 20°C
(D022+D023) [2003Kar]
in the as cast (Ti0.92Zr0.08)Al3 alloy
(D022+D023) [2003Kar]
in the as cast (Ti0.96Zr0.04)Al3 alloy
(D022+D023) [2003Kar]
in the (Ti0.84Zr0.16)Al3 alloy at 600°C
(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
Landolt-BörnsteinNew Series IV/11A4
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
Al 60.00Data / Grid: at.%
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
Al 60.00Data / Grid: at.%
Axes: at.%
γ
β
α
γ+βγ+α
Fig. 2: Al-Ti-Zr.
Partial isothermal
section at 1200°C
66
Landolt-BörnsteinNew Series IV/11A4
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
Al 60.00Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
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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].
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After prolonged annealing at 1200°C the X-ray diffractograms revealed TiB2 as a secondary phase
[1975Aiv].
Invariant Equilibria
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”.
Temperature – Composition Sections
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
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[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
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concentration sections TiB2-C, TiB2-TiC1-x and TiB2-’B4C’. These isopleths in the original version lack
information concerning the phase Ti3B4.
Notes on Materials Properties and Applications
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.
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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)
[1982McC] McCauley, J.W., Corbin, N.D., Resetar, T., Wong, P., “Simultaneous Preparation and
Self-Sintering of Materials in the System Ti-B-C”, Ceram. Eng. Sci. Proc., 3, 538-554
(1982) (Equi. Diagram, Experimental, 9)
[1982Unr] Unrod, V.I., “Correlation of Abrasive-Mechanical Properties with the Structure of Alloys in
the Titanium Carbide - Titanium Boride System” (in Russian), Vysoko-Temperaturnye
Boridy i Silitsidy, Kiev, 97-100 (1982) (Equi. Diagram, Experimental)
[1983Shc] Shcherbakov, V.A., Pityulin, A.N., “Reactions in Titanium-Carbon-Boron Mixtures” (in
Russian), Fiz. Goreniya Vzryva, 19, 108-111 (1983)
[1983Sco] Shouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties
of the Metal - Carbon - Nitrogen and Metal - Carbon - Boron Systems” (in French), Rev. Int.
Hautes Tempér., Refract. Fr., 20, 261-311 (1983) (Review, 154)
[1984Hol] Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in
German), Materialkundlich - Technische Reihe, Vol. 6 - Gebrüder Bornträger, Berlin, p.
264-274 (1984) (Crys. Structure, Equi. Diagram, Review, 87)
[1985Cor] Corbin, N.D., Resetar, T., McCauley, J.W., “‘SHS’ Self Sintering of Materials in the
Titanium - Boron - Carbon System”, Plenum Press., New York, USA, 337-346 (1985)
(Equi. Diagram, Experimental, 7)
[1985Nis] Nishiyama, K., Umekawa, S., “Boron Carbide - Titanium Diboride Composites”, Trans.
Jap. S. C. M., 11, 53-62 (1985) (Experimental, Crys. Structure, 6)
[1986Hal] Halverson, D.C., Munir, Z.A., “Boron Carbide Reactive Metal Cermets: Thermodynamic
Considerations in Boron Carbide-Titanium Cermets”, Ceram. Eng. Sci., Proc., 7,
1001-1010 (1986) (Thermodyn.)
[1986Mur] Murray, J.L., Liao, P.K., Spear, K.E., “The B-Ti (Boron - Titanium) System”, Bull. Alloy
Phase Diagrams, 7, 550-555 (1986) (Equi. Diagram, Thermodyn., Theory, Review, #, 48)
[1986Ord] Ordan'yan, S.S., Stepanenko, E.K., Dmitriev, A.I., Shchemeleva, M.V., “Interaction in the
Boron Carbide - Titanium Boride (B4C-TiB2) System” (in Russian), Sverkhtverd. Mater., 5,
27-29 (1986) (Equi. Diagram, Experimental, #, 4)
[1990Ase] Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral
Borides”, in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R.,
(Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U. K., Sept. 1989,
published as ASI-Series, Series E: Applied Sciences -Vol. 185, Kluwer Acad. Publ.,
Dordrecht 97-111 (1990) (Crys. Structure, Review, Experimental, 14)
[1990Kan] Kank, E.S., Kim, C.H., “Improvements in Mechanical Properties of TiB2 by the Dispersion
of B4C Particles”, J. Mater. Sci., 25, 580-584 (1990) (Experimental, 21)
[1990Kno] Knotek, O., Jungbluth, F., Breidenbach, R., “Magnetron Sputtered Superhard Coatings
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)
[1990Mit] Mitterer, C., Rauter, M., Rödhammer, P., “Sputter Deposition of Ultrahard Coatings within
the System Ti-B-C-N”, Surf. Coat. Technol., 41, 351-363 (1990) (Experimental)
[1990Spe] Spencer, P. J., Holleck, H., “Application of a Thermochemical Data Bank System to the
Calculation of Metastable Phase Formation during PVD of Carbide, Nitride, and Boride
Coating”, High Temp. Sci., 27, 295-309 (1990) (Equi. Diagram, Theory, #, 14)
[1990Sun] Sundman, B., “Review of Alloys Modelling”, An. Fis., Ser. B, 86, 69-82 (1990)
(Theory, 24)
75
Landolt-BörnsteinNew Series IV/11A4
MSIT®
B–C–Ti
[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)
[1991Hil] Hilz, G., Holleck, H., “Characterization of Microstructure and Interfaces in TiC-TiB2
Coatings”, Mat. Sci. Eng., A139, 268-275 (1991) (Experimental, 11)
[1991Hol] Holleck, H., Lahres, M., “Two-Phase TiC/TiB2 Hard Coatings”, Mat. Sci. Eng., A140,
609-615 (1991) (Experimental, Equi. Diagram, 19)
[1991Lev] Levashov, E.A., Borovinskaya, I.P., Bogatov, Yu.V., Kost, A.G., Milovidov, A.A.,
“Structure Formation in Self-Propagating High-Temperature Synthesis of Titanium
Carbide-Titanium Diboride Composites” (in Russian), Izv. Vyssh. Uchebn. Zaved., Chern.
Metall., (3) 82-86 (1991) (Experimental, 12)
[1991Pak] Pakholkov, V.V., Brettser-Portnov, I.V., Grigorov, I.G., Aliamovskii, S.I., Zainulin, Y.G.,
“Phase Formation in the System TiB-TiC” (in Russian), Zh. Neorg. Khim., 36(6),
1604-1609 (1991) (Experimental, Crys. Strucrture, 9)
[1992Ale] Alekseeva, Z., “Boron - Carbon - Titanium”, Leuven Proceeding of the COST 507, New
Ligh Alloys, Part B, Effenberg, G. (Ed.), Commission of the European Communities,
184-201 (1992) (Review, #, 25)
[1993Wer] Werheit, H., Kuhlmann, U., Laux, M., Lundström, T., “Structural and Electronic Properties
of Carbon-Doped -Rhombohedral Boron”, Phys. Stat. Sol., B179, 489-511 (1993) (Crys.
Structure, Experimental, 51)
[1994Bae] Baetzner, C., Thesis, Max-Planck-Institut - PML, Stuttgart, Germany (1994).
[1995Dus] Duschanek, H., Rogl, P., Lukas, H.L., “A Critical Assessment and Thermodynamic
Calculation of the Boron - Carbon - Titanium (B-C-Ti) Ternary System”, J. Phase Equilib.,
16(1), 46-60 (1995) (Thermodyn., Equi. Diagram, Review, 36)
[1995Gui] Guiban, M.A., Male, G., “Experimental Study of the Ti-B-C System Using LPCVD”,
J. Eur. Ceram. Soc., 15, 537-549 (1995) (Experimental, 12)
[1995Vil] Villars, P., Prince, A., Okamoto, H., Handbook of Ternary Alloys Phase Diagrams, Vol. 5,
ASM International, Materials Park, Ohio, USA (1995) (Equi. Diagam, Crys. Structure,
Review, 8)
[1996Bro] Brodkin, D., Barsoum, M.W., “Isothermal Section of Ti-B-C Phase Diagram at 1600°C”,
J. Am. Ceram. Soc., 79(3), 785-87 (1996) (Experimental, Equi. Diagram, 11)
[1996Kas] Kasper, B., Thesis, Max-Planck-Institut - PML, Stuttgart, Germany (1996).
[1997Dum] 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”, TRITA-MAC-0616,
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
Landolt-BörnsteinNew Series IV/11A4
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)
Table 1: Crystallographic Data of Solid Phases
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,
from alloy Ti25B15C60, quenched from
2838°C, contained also TiB2 and TiC1-x
[1966Rud]
from alloy Ti15B30C55, quenched from
2636°C, contained also TiB2 [1966Rud]
77
Landolt-BörnsteinNew Series IV/11A4
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,
quenched from 2400°C [1986Ord]
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]
from alloy Ti73B20C7, quenched from
1580°C, contained also ( Ti) and TiC1-x
[1966Rud]
from alloy Ti64B29C7, quenched from
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]
from alloy Ti49B46C5, annealed at 1550°C,
contained also TiC1-x and TiB [1995Dus]
from alloy Ti49B31C20, annealed at 1550°C,
contained also TiC1-x and TiB2 [1995Dus]
Phase/
Temperature Range
[°C]
Pearson Symbol/
Space Group/
Prototype
Lattice Parameters
[pm]
Comments/References
78
Landolt-BörnsteinNew Series IV/11A4
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]
from alloy Ti49B31C20, annealed at 1550°C,
contained also Ti2C and Ti3B4 [1995Dus]
from alloy Ti42B23C25, quenched from
2620°C, contained also TiC1-x [1966Rud]
from alloy Ti35B60C5, quenched from
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]
quenched from 2980°C [1986Ord]
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
Landolt-BörnsteinNew Series IV/11A4
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]
from alloy Ti49B46C5, annealed at 1550°C,
contained TiB and Ti3B4 [1995Dus]
from alloy Ti59B20C21, quenched from
2460°C, contained ( Ti) and TiB2 [1966Rud]
from alloy Ti55B15C30, quenched from
2642°C, contained TiB2 [1966Rud]
from alloy Ti55B30C15, quenched from
2518°C, contained TiB, TiB2 [1966Rud]
from alloy Ti43B47C10, quenched from
2630°C, contained TiB2 [1966Rud]
from alloy Ti47B15C38, quenched from
2661°C, contained TiB2 [1966Rud]
from alloy Ti63B3C31, quenched from
2373°C, contained traces ( Ti) [1966Rud]
from alloy Ti62B3C35, quenched from
2800°C, [1966Rud]
from alloy Ti57B3C40, quenched from
2800°C, [1966Rud]
from alloy Ti52B3C45, quenched from
2992°C, [1966Rud]
from alloy Ti33B21C46, quenched from
2517°C, contained TiB2 and B4C [1966Rud]
from alloy Ti25B30C45, quenched from
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Calculated isothermal
section at 1400°C
Fig. 6: B-C-Ti.
Experimentally
derived isopleth
TiB2 - 'B4C'; after
[1966Rud]
86
Landolt-BörnsteinNew Series IV/11A4
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.
Calculated isothermal
section at 2150°C
Fig. 9: B-C-Ti.
Calculated isothermal
section at 2300°C
87
Landolt-BörnsteinNew Series IV/11A4
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.
Calculated isothermal
section at 2400°C
Fig. 11: B-C-Ti.
Calculated isopleth
Ti0.5B0.5 - C
88
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
B–Li–N
Pseudobinary Systems
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),
809-811 (1961) (Crys. Structure, Experimental, 4)
[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
Landolt-BörnsteinNew Series IV/11A4
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,
3547-3556 (1993) (Equi. Diagram, Experimental, 35)
[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
(1993) (Equi. Diagram, Experimental, 26)
[1993Wer] Werheit, H., Kuhlmann, U., Laux, M., Lundström, T., “Structural and Electronic Properties
of Carbon-Doped -Rhombohedral Boron”, Phys. Status Solidi (B), B179, 489-511 (1993)
(Crys. Structure, Experimental, 51)
[1995Pav] Pavlyuk, V., Bodak, O., “Boron-Lithium-Nitrogen”, MSIT Ternary Evaluation Program, in
MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services
GmbH, Stuttgart; Document ID: 10.20619.1.20, (1995) (Crys. Structure, Equi. Diagram,
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
MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services,
GmbH, Stuttgart, to be published, (2002) (Equi. Diagram, Review, 50)
Table 1: Crystallographic Data of Solid Phases
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]
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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
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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]
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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
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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
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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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B–Mg–N
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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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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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.
Invariant Equilibria
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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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].
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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.
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(Experimental, Crys. Structure, Equi. Diagram, 14)
[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
Landolt-BörnsteinNew Series IV/11A4
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,
Crys. Structure, Equi. Diagram, 38)
[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)
(Experimental, Crys. Structure, 17)
[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.
Structure, Theory, 27)
[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
Landolt-BörnsteinNew Series IV/11A4
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B–Mg–N
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
B–N–Ti
Temperature – Composition Sections
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].
Notes on Materials Properties and Applications
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),
1739-1742 (1971) (Experimental, Equi. Diagram, 9)
[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
Landolt-BörnsteinNew Series IV/11A4
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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)
(Experimental, Equi. Diagram, 20)
[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)
[1993Wer] Werheit, H., Kuhlmann, U., Laux, M., Lundström, T., “Structural and Electronic Properties
of Carbon-Doped -Rhombohedral Boron”, Phys. Status Solidi (B), B179, 489-511 (1993)
(Crys. Structure, Experimental, 51)
[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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
B–N–Ti
GmbH, Stuttgart; Document ID: 10.12201.1.20, (1996) (Crys. Structure, Equi. Diagram,
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)
(Experimental, Equi. Diagram, 19)
[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)
(Experimental, Equi. Diagram, 15)
[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,
201-209 (2001) (Review, Equi. Diagram, 65)
[2003Rec] Record, M.Ch., Tedenac, J.-C., “B-N (Boron-Nitrogen)”, MSIT Evaluation Program, in
MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services,
GmbH, Stuttgart, to be published, (2002) (Review, Equi. Diagram, 50)
Table 1: Crystallographic Data of Solid Phases
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]
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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
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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.
Isothermal section at
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.
Isothermal section at
1500°C under 105 Pa
of argon (in the
absence of external
nitrogen)
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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
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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].
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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.
Pseudobinary Systems
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.
Invariant Equilibria
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.
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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.
Notes on Materials Properties and Applications
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
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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
(1950) (Experimental, Crys. Structure, 14)
[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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
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
(1987) (Crys. Structure, Experimental, 10)
[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.
Structure, Experimental, 13)
[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.,
126
Landolt-BörnsteinNew Series IV/11A4
MSIT®
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.
Structure, Experimental, 20)
[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)
127
Landolt-BörnsteinNew Series IV/11A4
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C–N–Ti
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
C–N–Ti
Table 2: Invariant Equilibria
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
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
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.
Isothermal section at
500°C according to
[1978Arb1]
131
Landolt-BörnsteinNew Series IV/11A4
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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
N 50.00Data / Grid: at.%
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
N 50.00Data / Grid: at.%
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.
Isothermal section at
the invariant
temperature 1820°C,
calculated by
[1996Jon]
Fig. 5: C-N-Ti.
Tentative isothermal
section at 1150°C
132
Landolt-BörnsteinNew Series IV/11A4
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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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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.
Invariant Equilibria
[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.
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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.
Temperature – Composition Sections
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].
Notes on Materials Properties and Applications
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
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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.
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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)
(Experimental, Crys. Structure, 18)
[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)
[2002Tan1] Tang, K., Wang, C.A., Huang, Y., Zan, Q.F., Xu, X.L., “A Study on the Reaction
Mechanism and Growth of Ti3SiC2 Synthesized by Hot-Pressing”, Mater. Sci. Eng. A, 328,
206-212 (2002) (Theory, Crys. Structure, 22)
[2002Tan2] Tang, K., Wang, C.A., Xu, X., Huang, Y., “A Study of Powder X-ray Diffraction of
Ti3SiC2”, Mater. Lett., 55, 50-55 (2002) (Experimental, Crys. Structure, 12)
[2002Yok] Yokokawa, H., Yamauchi, S., Matsumoto, T., “Thermodynamic Database MALT for
Windows with Gem and CHD”, Calphad, 26(2), 155-166 (2002) (Calculation, Equi.
Diagram, 9)
[2002Yu] Yu, R., Zhan, Q., He, L.L., Zhou, Y.C., Ye, H.Q., “Polymorphism of Ti3SiC2”, J. Mater.
Res, 17(5), 948-950 (2002) (Experimental, Crys. Structure, 18)
[2003Jor] Jordan, J.L., Sekine, T., Kobayashi, T., Li, X., Thadhani, T., El-Raghy, T., Barsoum, M.V.,
“High Pressure Behavior of Titanium-Silicon Carbide (Ti3SiC2)”, J. Appl. Phys., 93,
9639-9643 (2003) (Experimental, Equi. Diagram, #, 8)
140
Landolt-BörnsteinNew Series IV/11A4
MSIT®
C–Si–Ti
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
C–Si–Ti
Table 2: Invariant Equilibria
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]
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
C–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti C
Si Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
C–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti C
Si Data / Grid: at.%
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
Si Data / Grid: at.%
Axes: at.%
(βTi)
τ1
SiC
TiSi2
TiSi
Ti5Si
4
T2
TiCx
(Si)
L
L
Fig. 7: C-Si-Ti.
Isothermal section at
1250°C [2000Du]
Fig. 6: C-Si-Ti.
Isothermal section at
1400°C [2000Du]
147
Landolt-BörnsteinNew Series IV/11A4
MSIT®
C–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti C
Si Data / Grid: at.%
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.
Isothermal section at
1100°C [2000Du]
Fig. 9: C-Si-Ti.
Isopleth at 5 at.% C
[2000Du]
148
Landolt-BörnsteinNew Series IV/11A4
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.
Isopleth at 10 at.% C
[2000Du]
Fig. 10b: C-Si-Ti.
Isopleth at 10 at.% C
from 33 to 38 at.% Si
[2000Du]
149
Landolt-BörnsteinNew Series IV/11A4
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.
Isopleth at 15 at.% C
[2000Du]
Fig. 12: C-Si-Ti.
The region of
temperature-pressure
stability range of the
Ti3SiC0.4 phase
[1993Sam]
150
Landolt-BörnsteinNew Series IV/11A4
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
lattice parameters of
Ti5Si3Cx (0< x<1)
Fig. 14: C-Si-Ti.
Relationship between
the deposited phase
and the deposition
condition [1987Got]
151
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
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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].
Invariant Equilibria
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
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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.
Structure, Experimental, 17)
[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)
(Equi. Diagram, Experimental, 14)
[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,
116-123 (1996) (Crys. Structure, Experimental, 13)
[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
Landolt-BörnsteinNew Series IV/11A4
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Ce–Mg–Y
Table 1: Crystallographic Data of Solid Phases
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]
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Landolt-BörnsteinNew Series IV/11A4
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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
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Ce–Mg–Y
Table 2: Invariant Equilibria
Reaction T [°C] Type Phase Composition (at.%)
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
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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.
Isothermal section at
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
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Landolt-BörnsteinNew Series IV/11A4
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Ce–Mg–Y
10
10
90
Y 15.00
Ce 0.00
Mg 85.00
Y 0.00
Ce 15.00
Mg 85.00
Mg Data / Grid: at.%
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
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Landolt-BörnsteinNew Series IV/11A4
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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.
Partial isopleth at 15
mass% Ce, from 0 to
20 mass% Y
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Landolt-BörnsteinNew Series IV/11A4
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.
Invariant Equilibria
[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.
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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.
Temperature – Composition Sections
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)
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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)
(Crys. Structure, Experimental, 0)
[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)
Table 1: Crystallographic Data of Solid Phases
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]
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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
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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
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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
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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.
Polythermal section at
34 mass% Zn
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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
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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].
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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.
Pseudobinary Systems
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.
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Co–Ni–Ti
Notes on Materials Properties and Applications
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),
(Experimental, Equi. Diagram, 5)
[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
Landolt-BörnsteinNew Series IV/11A4
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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)
(Experimental, Thermodyn., 2)
[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.
Structure, Theory, 101)
[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
Landolt-BörnsteinNew Series IV/11A4
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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)
Table 1: Crystallographic Data of Solid Phases
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]
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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
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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
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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
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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.
Isothermal section at
900°C
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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
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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.
Pseudobinary Systems
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].
Notes on Materials Properties and Applications
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
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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
(1962) (Crys. Structure, Experimental, 5)
[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)
(Crys. Structure, Experimental, 27)
[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)
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Landolt-BörnsteinNew Series IV/11A4
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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)
Table 1: Crystallographic Data of Solid Phases
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]
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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]
dissolves ~3 at.% Si
[1966Mar1]
Ti2Co (c)
< 1058
cF96
Fd3m
NiTi2
a = 1129.5 66.7 to 67.1 at.% Ti
[Mas2, V-C2]
dissolves ~5 at.% Si
[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]
dissolves ~2 at.% Co
[1966Mar1]
TiSi2< 1488
oF24
Fddd
TiSi2
a = 826.71 0.09
b = 480.00 0.05
c = 855.05 0.11
[Mas2, V-C2]
dissolves ~2 at.% Co
[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
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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
Si Data / Grid: at.%
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.
Isothermal section at
800°C
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Co–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Co
Si Data / Grid: at.%
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.
Isothermal section at
1100°C
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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
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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.
Liquidus and Solidus Surfaces
[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.
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Cr–Nb–Ti
Temperature – Composition Sections
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.
Notes on Materials Properties and Applications
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.
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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,
Experimental, #, *, 8)
[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
(1997) (Experimental, Mechan. Prop., 20)
[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
(1998) (Crys. Structure, Experimental, 25)
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Landolt-BörnsteinNew Series IV/11A4
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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
MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services,
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)
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Cr–Nb–Ti
Table 1: Crystallographic Data of Solid Phases
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
pure Ti at 25°C [Mas2]
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
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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.
Isothermal section at
1900°C
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
Axes: at.%
(Nb)
L+(Nb)
L
L+(Cr)
(Cr)Fig. 2: Cr-Nb-Ti.
Isothermal section at
1800°C
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Cr–Nb–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
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
Cr Data / Grid: at.%
Axes: at.%(Cr)
L+(Cr)+βNbCr2(h)
βNbCr2(h)
αNbCr2(r)
L+β+αNbCr2(r)L
(βTi) (Nb)
β
Fig. 3: Cr-Nb-Ti.
Isothermal section at
1700°C
Fig. 4: Cr-Nb-Ti.
Isothermal section at
1600°C
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Cr–Nb–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
Axes: at.%
β
(βTi) (Nb)
(Cr)
αNbCr2(r)
βTiCr2(h)
β+βTiCr2(h)+αNbCr
2(r)
β+βTiCr2(h)+αNbCr
2(r)
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
Axes: at.%
β
(Nb)
β+βTiCr2(h)+αNbCr
2(r)
L+β+βTiCr2(h)
L+(Cr)+βTiCr2(h)
βTiCr2(h)
(βTi)
αNbCr2(r)
L
(Cr)
β+βTiCr2(h)+αNbCr
2(r)
Fig. 6: Cr-Nb-Ti.
Isothermal section at
1400°C
Fig. 5: Cr-Nb-Ti.
Isothermal section at
1500°C
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Cr–Nb–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
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
Cr Data / Grid: at.%
Axes: at.%
(βTi) (Nb)
(Cr)+αNbCr2(r)
αNbCr2(r)
βTiCr2(h)
(Cr)
β
(Cr)+βTiCr2(h)+αNbCr
2(r)
β+βTiCr2(h)+αNbCr
2(r)
Fig. 7: Cr-Nb-Ti.
Isothermal section at
1300°C
Fig. 8: Cr-Nb-Ti.
Isothermal section at
1250°C
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Cr–Nb–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
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
Cr Data / Grid: at.%
Axes: at.%(Cr)
βTiCr2(h)
αNb1-x
TixCr
2(r)
β
(βTi) (Nb)
Fig. 9: Cr-Nb-Ti.
Isothermal section at
1200°C
Fig. 10: Cr-Nb-Ti.
Isothermal section at
1150°C
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Cr–Nb–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
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.
Isothermal section at
950°C
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
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.
Isothermal section at
1000°C
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Cr–Nb–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Nb
Cr Data / Grid: at.%
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
Cr Data / Grid: at.%
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.
Isothermal section at
800°C
Fig. 14: Cr-Nb-Ti.
Partial isothermal
section at 600°C
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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
constant mass ratio of
Nb:Ti=1:4
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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)+αNb1-xTixCr2(r)
(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
constant mass ratio of
Nb:Ti=2:3
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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)
(βTi,Nb)+αNb1-xTixCr2(r)
Fig. 19: Cr-Nb-Ti.
An isopleth at a
constant atomic ratio
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)
(βTi,Nb)+αNb1-xTixCr2(r)
Fig. 20: Cr-Nb-Ti.
An isopleth at a
constant atomic ratio
of Nb:Ti=3:1
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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%
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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
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Landolt-BörnsteinNew Series IV/11A4
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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.
Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
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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.
Temperature – Composition Sections
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.
Notes on Materials Properties and Applications
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,
Crys. Structure, 37)
[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)
(Equi. Diagram, Experimental, 8)
[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.
Diagram, Experimental, 5)
[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
Landolt-BörnsteinNew Series IV/11A4
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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
Titanium System”, ASM Internal, Metals Park, OH, 197-210 (1987) (Equi. Diagram, Crys.
Structure, Thermodyn., Review, 110)
[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)
[1992Shi] Shimizu, K., Tadaki, T., “Recent Studies on the Precise Crystal-Structural Analyses of
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)
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Cr–Ni–Ti
[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)
Table 1: Crystallographic Data of Solid Phases
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]
pure Ti 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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Table 2: Invariant Equilibria
Reaction T [°C] Type Phase Composition (at.%)
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%
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20
40
60
80
20 40 60 80
20
40
60
80
Ti Cr
Ni Data / Grid: at.%
Axes: at.%
(Ni)
(βCr)
TiNi3
TiNi
Ti2Ni
(βTi)
(αTi) αTiCr2
βTiCr2
Fig. 2: Cr-Ni-Ti.
Isothermal section at
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.
Partial isopleth at 20
mass% Cr
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Cr–Ni–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Cr
Ni Data / Grid: at.%
Axes: at.%
(Ni)
(βCr)
TiNi3
TiNi
L
(βTi)
αTiCr2
βTiCr2
Fig. 3: Cr-Ni-Ti.
Isothermal section at
1027°C
20
40
60
80
20 40 60 80
20
40
60
80
Ti Cr
Ni Data / Grid: at.%
Axes: at.%
(βCr)
(Ni)
TiNi3
NiTi L
(βTi)
γTiCr2
Fig. 4: Cr-Ni-Ti.
Isothermal section at
1277°C
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20
40
60
80
20 40 60 80
20
40
60
80
Ti Cr
Ni Data / Grid: at.%
Axes: at.%
(Ni)
(βCr)(βTi)
TiNi3
L
γTiCr2
Fig. 5: Cr-Ni-Ti.
Isothermal section at
1352°C
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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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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.
Pseudobinary Systems
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.
Invariant Equilibria
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
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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.
Temperature – Composition Sections
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
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Cu–Mg–Ni
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.
Structure, Experimental, 15)
216
Landolt-BörnsteinNew Series IV/11A4
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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,
513-519 (1979) (Review, Equi. Diagram, 7)
[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
Landolt-BörnsteinNew Series IV/11A4
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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)
Table 1: Crystallographic Data of Solid Phases
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]
at x = 0 [Mas2, 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]
at x = 0 [Mas2, V-C]
* (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
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Cu–Mg–Ni
Table 2: Invariant Equilibria
* 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.
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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
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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
Mg Data / Grid: at.%
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
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Cu–Mg–Ni
20
40
60
80
20 40 60 80
20
40
60
80
Ni Cu
Mg Data / Grid: at.%
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
Mg Data / Grid: at.%
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.
Calculated isothermal
section at 475°C
222
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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
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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
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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
Landolt-BörnsteinNew Series IV/11A4
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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.
Pseudobinary Systems
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.
Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
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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].
Temperature – Composition Sections
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,
1603-1605 (1933) (Equi. Diagram, Experimental, #, 2)
227
Landolt-BörnsteinNew Series IV/11A4
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
(1957) (Experimental, 28)
[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
Landolt-BörnsteinNew Series IV/11A4
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,
325-328 (1984) (Crys. Structure, Experimental, 4)
[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),
1668-1670 (1984) (Crys. Structure, Experimental, 13)
[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
Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science
International Services GmbH, Stuttgart; Document ID: 20.12505.1.20, (2002) (Crys.
Structure, Equi. Diagram, Assessment, 23)
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Mg–Si
Table 2: Invariant Equilibria
', 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]
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Reaction T [°C] Type Phase Composition (at.%)
Cu Mg Si
232
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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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
Si 35.00Data / Grid: at.%
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
Si Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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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
Si 50.00Data / Grid: at.%
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
Si Data / Grid: at.%
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.
Isothermal section at
500°C in the 3 region
Fig. 5: Cu-Mg-Si.
Isothermal section at
450°C
236
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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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].
Pseudobinary Systems
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.
Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Mg–Sn
Temperature – Composition Sections
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
Landolt-BörnsteinNew Series IV/11A4
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)
(Experimental, Crys. Structure, 9)
[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,
93, 323-330 (1987) (Experimental, Crys. Structure, 16)
[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.
Diagram, Experimental, 4)
[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
Landolt-BörnsteinNew Series IV/11A4
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Cu–Mg–Sn
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Mg–Sn
Table 2: Invariant Equilibria
', 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]
Reaction T [°C] Type Phase Composition (at.%)
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
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Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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
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Landolt-BörnsteinNew Series IV/11A4
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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
Sn Data / Grid: at.%
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.
Isothermal section at
400°C
248
Landolt-BörnsteinNew Series IV/11A4
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].
Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
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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),
1252-1255 (1966) (Equi. Diagram, Experimental, #, *, 9)
[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)
Table 1: Crystallographic Data of Solid Phases
Table 2: Invariant Equilibria
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]
Reaction T [°C] Type Phase Composition (at.%)
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
-
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Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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]
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Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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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].
Pseudobinary Systems
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.
Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
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Cu–Mg–Zn
Temperature – Composition Sections
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.
Structure, Experimental, 10)
[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
Landolt-BörnsteinNew Series IV/11A4
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),
76-93 (1950) (Equi. Diagram, Experimental, #, *, 13)
[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,
396-401 (1952) (Equi. Diagram, Experimental, #, *, 17)
[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
(1958) (Experimental, 5)
[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
Landolt-BörnsteinNew Series IV/11A4
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)
(Crys. Structure, Experimental, 13)
[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.
Structure, Experimental, 13)
[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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Mg–Zn
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Mg–Zn
Table 2: Invariant Equilibria
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]
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
lattice parameters of
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
Landolt-BörnsteinNew Series IV/11A4
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
dataset of [1998Lia]
261
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Zn Data / Grid: at.%
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
dataset of [1998Lia]
Fig. 6: Cu-Mg-Zn.
Zn corner of the
liquidus surface,
calculated from the
dataset of [1998Lia]
264
Landolt-BörnsteinNew Series IV/11A4
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
dataset of [1998Lia]
Fig. 8: Cu-Mg-Zn.
Solvus surface of the
(Zn) solid solution,
calculated from the
dataset of [1998Lia]
265
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Mg–Zn
20
40
60
80
20 40 60 80
20
40
60
80
Mg Cu
Zn Data / Grid: at.%
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.
Isothermal section at
340°C, calculated
from the dataset of
[1998Lia]
266
Landolt-BörnsteinNew Series IV/11A4
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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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].
Pseudobinary Systems
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].
Invariant Equilibria
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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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.
Temperature – Composition Sections
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].
Notes on Materials Properties and Applications
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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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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System”, Russ. Metall., (2), 192-195 (1992), translated from Izv. Akad. Nauk SSSR, Met.,
(2) 223-226 (1992) (Equi. Diagram, Experimental, *, #, 9)
[1992Ali2] Alisova, S.P., Budberg, P.B., “Phase Diagram of the Ti-Cu-Ni System”, Russ. Metall., (4)
206-210 (1992) translated from Izv. Akad.Nauk SSSR,Met., (4) 218-223 (1992) (Equi.
Diagram, Experimental, 14)
[1992Ali3] Alisova, S.P., Budberg, P.B., Kovneristyi, Yu.K., “Phase Transformations in Copper Rich
Ti-Cu-Ni Alloys”, Russ. Metall., (5), 115-118 (1992) translated from Izv. Akad. Nauk SSSR,
Met., (5), 126-128 (1992) (Equi. Diagram, Experimental, *, 6)
[1992Ali4] Alisova, S.P., Kovneristyi, Yu.K., Budberg, P.B., “Phase Equilibria in the Titanium Based
Secondary Systems Based on Intermetallic Compounds of Titanium and Some Peculiarities
of the Diffusionless Solidifying of the Alloys” (in Russian), Metalloved. i Obrab. Tsvet.
Splavov, Nauka, Moscow, 17-24 (1992) (Equi. Diagram, Experimental, 4)
[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. Met.
Mater., 26, 1163-1168 (1992) (Crys. Structure, Experimental, 13)
[1992Yu] Yu, W., “The Application of Thermal Analysis in the Study of Ni-Ti Shape Memory
Alloys” in “Thermal Analysis in Metallurgy”, Shull, R.D., Joshi, A. (Eds.), The Minerals,
Metals & Materials Soc., Warrendale, PA, USA, 187-201 (1992) (Crys. Structure,
Experimental, 34)
[1993Ali1] Alisova, S.P., Budberg, P.B., Barmina, T.I., Lutskaya, N.V., “Metastable Phase Diagram of
Ti-Cu-Ni System”, Russ. Metall., (2), 177-180 (1993) translated from Izv. Ross. Akad.
Nauk, Met., (2), 205-210 (1993) (Equi. Diagram, Experimental, 7)
[1993Ali2] Alisova, S.P., Lutskaya, N.V., Budberg, P.B., Bychkova, E.I., “Phase Constitution of the
TiCu-TiNi-TiCo (TiFe) Systems in the Equilibrium and Metastable States”, Russ. Metall.,
(3), 205-212 (1993) translated from Izv. Ross. Akad. Nauk, Met., 3, 222-229, (1993) (Equi.
Diagram, Experimental, 16)
[1993Lo] Lo, Y.C., Wu, S.K., Horng, H.E., “A Study of B2 - B19 - B19' Two-Stage Martensitic
Transformation in a Ti50Ni40Cu10 Alloy”, Acta Metall. Mat., 41(3), 747-759 (1993)
(Experimental, 43)
[1993Mat1] Matveeva, N.M., Bashanova, N.N., Lovtsova, I.D., “Some Characteristics of Shape
Memory Efeect and Mechanical Properties of Ti50Ni25Cu25 Alloy” (in Russian), Metally,
4, 197-199 (1993) (Experimental, 5)
[1993Mat2] 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. Ross. Akad. Nauk, (3), 233-237 (1993) (Theory, 10)
[1994Jia] Jia, C.C., Ishida, K., Nishizawa, T., “Partitioning of Alloying Elements Between gamma
(A1) and eta (DO24) Phases in The Ni-Ti Base Systems”, Exp. Methods Phase Diagram
Determ., Proc. Symp., 1993, 31-38, (1994) (Experimental, 8)
[1996Fuk] Fukuda, T., Kitayama, M., Kakeshita, T., Saburi, T., “Martensitic Transformation Behavior
of a Shape Memory Ti-40.5Ni-10Cu Alloy. Affected by the C11b-Type Precipitates”,
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
(1997) (Experimental, 14)
[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
Landolt-BörnsteinNew Series IV/11A4
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
(1997) (Experimental, 11)
[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
Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science
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
Landolt-BörnsteinNew Series IV/11A4
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
Reference Experimental Techniques Temperature/Composition/Phase Range
275
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Ni–Ti
Table 2: Crystallographic Data of Solid Phases
[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
[1997Pus3] XRD, Electron diffraction, TEM,
Electrical resistivity
Ti50Ni50-xCux 0 < x < 40
[1997Pus4] XRD, Electron diffraction, TEM,
Electrical resistivity
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]
Reference Experimental Techniques Temperature/Composition/Phase Range
276
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Ni–Ti
Table 3: Invariant Equilibria
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Cu Data / Grid: at.%
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.
Isothermal section at
870°C [1978Loo]
Fig. 4: Cu-Ni-Ti.
Partial liquidus
surface
282
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Ni–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Ni
Cu Data / Grid: at.%
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.
Isothermal section at
800°C [1978Loo]
Fig. 7: Cu-Ni-Ti.
Polythermal section
Cu-Ni3Ti [1967Jos2]
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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
Landolt-BörnsteinNew Series IV/11A4
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
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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).
Pseudobinary Systems
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.
Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
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Cu–Si–Ti
Temperature – Composition Sections
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.
Notes on Materials Properties and Applications
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,
Experimental, #, *, 79)
[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”,
665-668 (1979) (Review, Equi. Diagram, 7)
[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
Landolt-BörnsteinNew Series IV/11A4
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.
Diagram, Experimental, 7)
[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)
[1996Sei] Seifert, H.J., Lukas, H.L., Petzow, G., “Thermodynamic Optimization of the Ti-Si System”,
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
Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science
International Services GmbH, Stuttgart; Document ID: 20.12505.1.20, (2002) (Crys.
Structure, Equi. Diagram, Assessment, 23)
Table 1: Crystallographic Data of Solid Phases
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]
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Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Si–Ti
1) Ti2Cu is likely to exhibit a small homogeneity region (see also sections “Solid Phases” and “Isothermal
Sections”).
Table 2: Invariant Equilibria
*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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Cu–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Cu
Si Data / Grid: at.%
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.
Isothermal section at
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
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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
Si 40.00Data / Grid: at.%
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]
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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
Si 40.00Data / Grid: at.%
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
Si 40.00Data / Grid: at.%
Axes: at.%
Ti5Si
3(ζ)
(βTi)
Ti2Cu(ε)
γ
(αTi)
γ+ζ+ε
(αTi)+γ+ε
(αTi)+(βTi)+ε
Fig. 6: Cu-Si-Ti.
Isothermal section at
830°C; after
[1976Ali]
Fig. 5: Cu-Si-Ti.
Isothermal section at
810°C; after
[1976Ali]
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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
Si 40.00Data / Grid: at.%
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
Si 40.00Data / Grid: at.%
Axes: at.%Ti
5Si
3(ζ)
γ
(βTi)
Ti2Cu(ε)
(βTi)+γ+ε
γ + ζ + ε
Fig. 7: Cu-Si-Ti.
Isothermal section at
870°C; after
[1976Ali]
Fig. 8: Cu-Si-Ti.
Isothermal section at
900°C; after
[1976Ali]
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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
Si 40.00Data / Grid: at.%
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
Si 40.00Data / Grid: at.%
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.
Isothermal section at
950°C; after
[1976Ali]
Fig. 10: Cu-Si-Ti.
Isothermal section
system at 1010°C;
after [1976Ali]
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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
Si 40.00Data / Grid: at.%
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
Si 40.00Data / Grid: at.%
Axes: at.%
(βTi)L
Ti5Si
3(ζ)
Ti2Cu(ε)
L+ε
L+(βTi)+ζ+ε
(βTi)+L
Fig. 11: Cu-Si-Ti.
Isothermal section at
1100°C; after
[1976Ali]
Fig. 12: Cu-Si-Ti.
Isothermal section at
1200°C; after
[1976Ali]
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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
Si 40.00Data / Grid: at.%
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
Si 60.00Data / Grid: at.%
Axes: at.%
L+ζ+ε+TiCu
L+ε+TiCu
L+ε+TiCuL
Ti5Si
3(ζ)
Ti2Cu TiCu
Fig. 13: Cu-Si-Ti.
Isothermal section at
1300°C; after
[1976Ali]
Fig. 14: Cu-Si-Ti.
Partial isothermal
section at 935°C; after
[1985Ali]
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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
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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.
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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.
Pseudobinary Systems
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.
Invariant Equilibria
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
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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].
Temperature – Composition Sections
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.
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Notes on Materials Properties and Applications
[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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Landolt-BörnsteinNew Series IV/11A4
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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.
304
Landolt-BörnsteinNew Series IV/11A4
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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)
305
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Fe–Ni–Ti
Table 1: Crystallographic Data of Solid Phases
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
pure Ti at 25°C [Mas2]
'-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
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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
Phase/ Temperature Range [°C]
Pearson Symbol/ Group Space/Prototype
Lattice Parameters [pm]
Comments/References
307
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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
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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
Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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Fe–Ni–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Fe
Ni Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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Fe–Ni–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Fe
Ni Data / Grid: at.%
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
Ni Data / Grid: at.%
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.
Isothermal section at 1000°C
312
Landolt-BörnsteinNew Series IV/11A4
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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.
Isothermal section at 700°C
Fig. 8: Fe-Ni-Ti.
A polythermal section at a constant Ti content of 8 mass%
313
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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
Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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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)
(αTi)+(βTi)+Ti(Fe,Ni)
(αTi)+(βTi)+Ti(Fe,Ni)
Fig. 15: Fe-Ni-Ti.
Partial polythermal section along the line of constant Fe:Ni atomic ratio of 3:1
317
Landolt-BörnsteinNew Series IV/11A4
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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].
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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.
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Pseudobinary Systems
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.
Notes on Materials Properties and Applications
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
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[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.
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[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,
98-103 (2000) (Crys. Structure, Experimental, 14)
[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)
(Crys. Structure, Experimental, 17)
[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.
Diagram, Review, #, 4)
[2001Yin] Yin, J., Yamada, T., Yoshinari, O., Tanaka, K., “Improvement of Hydrogen Storage
Properties of Mg-Ni Alloys by Rare-Earth Addition”, Mater. Trans., JIM, 42(4), 712-716
(2001) (Experimental, 12)
[2002Blo] Blomqvist, H., Noreus, D., “Mechanically Reversible Conductor-Insulator Transition in
Mg2NiH4”, J. Appl. Phys., 91(8), 5141-5148 (2002) (Crys. Structure, Electr. Prop.,
Experimental, 34)
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H–Mg–Ni
[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)
[2002Kuj] Kuji, T., Nakano, H., Aizawa, T., “Hydrogen Absorption and Electrochemical Properties of
Mg2-xNi (x=0-0.5) Alloys Prepared by Bulk Mechanical Alloying”, J. Alloy Compd.,
330/332, 590-596 (2002) (Crys. Structure, Experimental, #, 15)
[2002Li] Li, L., Saita, I., Saito, K., Akiyama, T., “Effect of Synthesis Temperature on the Activity of
Products in Hydriding Combustion Synthesis of Mg2NiH4”, J. Alloys Compd., 345, 189-195
(2002) (Experimental, Kinetics, 16)
[2002Mye] Myers, W.R., Wang, L.-W., Richardson, T.J., Rubin, M.D., “Calculation of
Thermodynamic, Electronic, and Optical Properties of Monoclinic Mg2NiH4”, J. Appl.
Phys., 91(8), 4879-4885 (2002) (Crys. Structure, Optical Prop., Thermodyn., 31)
[2002Par] Parker, S.F., Williams, K.P.J., Smith, T., Bortz, M., Bertheville, B., Yvon, K., “Vibration
Spectroscopy of Tetrahedral Ternary Metal Hydrides: Mg2NiH4, Rb3ZnH5 and their
Deuterides”, Phys. Chem. Chem. Phys., 4, 1732-1737 (2002) (Crys. Structure,
Experimental, 23)
[2002Tak] Takamura, H., Kakuta, H., Kamegawa, A., Okada, M., “Crystal Structure of Novel Hydrides
in a Mg-Ni-H System Prepared under an Ultra High Pressure”, J. Alloys Compd., 330/332,
157-161 (2002) (Crys. Structure, Experimental, 10)
[2003Sai] Saita, I., Li, L., Saito, K., Akiyama, T., “Hydriding Combustion Synthesis of Mg2NiH4”,
J. Alloys Compd., 356/357, 490-493 (2003) (Experimental, Phase Relation, Kinetics, 12)
[2003Yam] Yamamoto, S., Fukai, Y., Roennebro, E., Chen, J., Sakai, T., “Structural Changes of
Mg2NiH4 under High Hydrogen Pressure”, J. Alloys Compd., 356/357, 697-700 (2003)
(Crys. Structure, Experimental, 15)
[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)
Table 1: Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
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]
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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]
Phase/ Temperature Range [°C]
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
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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.
Isothermal section at
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.
Calculated isothermal
section at 301°C and
3.47 bar [1999Zen]
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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
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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].
Invariant Equilibria
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.
Temperature – Composition Sections
A calculated partial vertical section at 10 at.% Si is shown in Fig. 7. Three additional vertical sections are
given in [2004Kev].
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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.
Notes on Materials Properties and Applications
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.
Structure, Experimental, 15)
[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).
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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),
5083-5085 (1996) (Crys. Structure, Experimental, 14)
[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.
Structure, Experimental, 14)
[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)
Table 1: Crystallographic Data of Solid Phases
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]
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Li–Mg–Si
Table 2: Invariant Equilibria
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]
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Li–Mg–Si
20
40
60
80
20 40 60 80
20
40
60
80
Li Mg
Si Data / Grid: at.%
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
Si Data / Grid: at.%
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.
Calculated isothermal
section at 600°C
334
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Li–Mg–Si
20
40
60
80
20 40 60 80
20
40
60
80
Li Mg
Si Data / Grid: at.%
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
Si Data / Grid: at.%
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.
Calculated isothermal
section at 200°C
Fig. 5: Li-Mg-Si.
Calculated isothermal
section at 400°C
335
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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Li–Mg–Zn
Invariant Equilibria
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.
Structure, Experimental, 7)
[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.
Structure, Theory, 5)
338
Landolt-BörnsteinNew Series IV/11A4
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,
Experimental, #, 15)
[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)
[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)
[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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Li–Mg–Zn
Table 1: Crystallographic Data of Solid Phases
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]
dissolves 75.5 at.% Mg at 588°C [Mas2]
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
pure Mg at 25°C [V-C2]
dissolves 17 at.% Li at 588°C [Mas2]
at 0 to 18.4 at.% Li [V-C2]
dissolves 2.4 at.% Zn at 340°C [Mas2]
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]
dissolves 0.4 at.% Mg at 364°C [Mas2]
dissolves 1 at.% Li at 403°C [Mas2]
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
Landolt-BörnsteinNew Series IV/11A4
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°
at 60 at.% Zn [Mas2, 1992Aga]
~1 at.% of solubility range [Mas2]
[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
Landolt-BörnsteinNew Series IV/11A4
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
Zn Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Li–Mg–Zn
20
40
60
80
20 40 60 80
20
40
60
80
Li Mg
Zn Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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].
Invariant Equilibria
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].
Temperature – Composition Sections
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.
Notes on Materials Properties and Applications
[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
Landolt-BörnsteinNew Series IV/11A4
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.
Diagram, Experimental, 6)
[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,
200-204 (1971) (Equi. Diagram, Experimental, 4)
[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)
Table 1: Crystallographic Data of Solid Phases
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
pure Mg at 25°C [V-C2]
[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
Landolt-BörnsteinNew Series IV/11A4
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
at 84 at.% Mg [V-C2]
at 87 at.% Mg [V-C2]
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
Landolt-BörnsteinNew Series IV/11A4
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
Mg Data / Grid: at.%
Axes: at.%
500
566°C
600
E1, 536
(Mg)
NdMg3
Y5Mg
24
545°C
Fig. 3: Mg-Nd-Y.
Partial liquidus
surface
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Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Nd–Y
Y 5.00
Nd 0.00
Mg 95.00
Y 0.00
Nd 5.00
Mg 95.00
Mg Data / Grid: at.%
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
Mg Data / Grid: at.%
Axes: at.%
(Mg)
(Mg) + NdMg12
(Mg) + NdMg12
+ Y5Mg
24+x
(Mg) + Y5Mg
24+x
Fig. 5: Mg-Nd-Y.
Isothermal section at
400°C
Fig. 4: Mg-Nd-Y.
Isothermal section at
300°C
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Landolt-BörnsteinNew Series IV/11A4
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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
Mg Data / Grid: at.%
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.
Isothermal section at
500°C
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Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Pseudobinary Systems
[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.
Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
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)
(Review, Equi. Diagram, 27)
[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)
Table 1: Crystallographic Data of Solid Phases
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]
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Landolt-BörnsteinNew Series IV/11A4
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
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Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Nd–Zr
20
40
60
80
20 40 60 80
20
40
60
80
Nd Zr
Mg Data / Grid: at.%
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
Mg Data / Grid: at.%
Axes: at.%
(Mg)
(Mg) + (αZr)
(Mg) + (αZr) + Nd5Mg
41
(Mg) + Nd5Mg
41
Fig. 3: Mg-Nd-Zr.
Partial isothermal
section at 500°C
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Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Nd–Zr
Nd 0.20
Zr 0.00
Mg 99.80
Nd 0.00
Zr 0.20
Mg 99.80
Mg Data / Grid: at.%
Axes: at.%
(Mg)
(Mg) + (αZr)(Mg) + Nd
5Mg
41
(Mg) + (αZr) + Nd5Mg
41
Fig. 4: Mg-Nd-Zr.
Partial isothermal
section at 200°C
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Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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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.
Pseudobinary Systems
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].
Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
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.
Temperature – Composition Sections
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)
[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)
[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)
[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
(1989) (Experimental, Equi. Diagram, 20)
[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)
359
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Sn–Zn
[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)
Table 1: Crystallographic Data of Solid Phases
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]
dissolves 2.4 at.% Zn at 340°C [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
pure Zn at 25°C [V-C2]
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]
~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°
60 at.% Zn [Mas2, 1992Aga]
~1 at.% of solubility range [Mas2]
[1975Yar, V-C2]
labelled as Mg4Zn7 in [V-C2]
360
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Sn–Zn
Table 2: Invariant Equilibria
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Sn–Zn
20
40
60
80
20 40 60 80
20
40
60
80
Mg Zn
Sn Data / Grid: at.%
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
Sn Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Sn–Zn
20
40
60
80
20 40 60 80
20
40
60
80
Mg Zn
Sn Data / Grid: at.%
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.
Isothermal section at
250°C
Fig. 7: Mg-Sn-Zn.
Polythermal section at
40 mass% Sn
365
Landolt-BörnsteinNew Series IV/11A4
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.
Polythermal section at
85 mass% Sn
Fig. 9: Mg-Sn-Zn.
Polythermal section at
10 mass% Sn
366
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
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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.
Pseudobinary Systems
[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.%.
Invariant Equilibria
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
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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.
Notes on Materials Properties and Applications
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)
[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)
[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)
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Landolt-BörnsteinNew Series IV/11A4
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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,
(1974) (Experimental, Equi. Diagram, 6)
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[1977Ray] Raynor, G.V., “Constitution of Ternary and Some More Complex Alloys of Mg”, Int.
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[1979Dob] Dobatkina, T.V., “Solid Solubility of Yttrium and Zinc in Magnesium”, Russ. Metall., 2,
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the Mg-Y-Zn System”, Russ. Metall., 1, 4179-182 (1979) (Experimental, Equi. Diagram,
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[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)
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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),
263-273 (1998) (Experimental, Crys. Structure, 9)
[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,
63-67 (1999) (Crys. Structure, Experimental, 18)
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Zn-Mg-Rare-Earth Quasicrystals”, Mater. Science Eng., 294-296, 29-32 (2000)
(Experimental, Crys. Structure, 17)
[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.
Structure, Experimental, 21)
[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.
Structure, Experimental, 22)
[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)
(Experimental, Mechan. Prop., 16)
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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,
323-326 (2000) (Crys. Structure, Experimental, 9)
[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
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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)
(Experimental, Equi. Diagram, 13)
[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)
(Experimental, Mechan. Prop., 16)
[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)
(Experimental, Equi. Diagram, 12)
[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)
(Experimental, Mechan. Prop., 34)
[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)
(Experimental, Mechan. Prop., 10)
[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)
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Mg–Y–Zn
Table 1: Crystallographic Data of Solid Phases
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
pure Zn at 25°C [V-C2]
dissolves 0.4 at.% Mg at 364°C [Mas2]
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
at 52 at.% Zn [Mas2, 1992Aga]
~1 at.% of solubility range [Mas2]
[1959Ray]
determined initially as
hexagonal by [1933Tar]
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Y–Zn
Table 2: Invariant Equilibria
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Zn 30.00Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Y–Zn
20
40
60
80
20 40 60 80
20
40
60
80
Y Mg
Zn Data / Grid: at.%
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
Zn 80.00Data / Grid: at.%
Axes: at.%
Y5Mg
24+x
τ1
τ4
τ2
MgZn
(Mg)
(Mg)+τ
2 +τ4
(Mg)+
MgZ
n+τ2
Fig. 5: Mg-Y-Zn.
Isothermal section at
600°C
Fig. 4: Mg-Y-Zn.
Partial isothermal
section at 300°C
378
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Y–Zn
20
40
60
80
20 40 60 80
20
40
60
80
Y Mg
Zn Data / Grid: at.%
Axes: at.%
Y2Zn
17
τ3
τ2'
τ2
MgZn2
L
20
40
60
80
20 40 60 80
20
40
60
80
Y Mg
Zn Data / Grid: at.%
Axes: at.%
Y2Zn
17
MgZn2
τ3
τ2'
τ2
L
(Mg)
Fig. 6: Mg-Y-Zn.
Isothermal section at
500°C
Fig. 7: Mg-Y-Zn.
Isothermal section at
427°C
379
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Invariant Equilibria
[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
Landolt-BörnsteinNew Series IV/11A4
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
(1968) (Equi. Diagram, Experimental, 4)
[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)
[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)
[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)
[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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mg–Zn–Zr
[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)
Table 1: Crystallographic Data of Solid Phases
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]
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
dissolves 0.4 at.% Mg at 364°C [Mas2]
pure Zn at 25°C [V-C2]
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
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°
at 60 at.% Zn [Mas2, 1992Aga]
~1 at.% of solubility range [Mas2]
[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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Isothermal section at
300°C
385
Landolt-BörnsteinNew Series IV/11A4
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
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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].
Pseudobinary Systems
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.
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[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.
Invariant Equilibria
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].
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Liquidus and Solidus Surfaces
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).
Temperature – Composition Sections
[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].
Notes on Materials Properties and Applications
[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.
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[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.
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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
(2001) (Equi. Diagram, Experimental, 18)
[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)
(Equi. Diagram, Experimental, 22)
[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)
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mo–Si–Ti
Table 2: Experimentally Determined Ternary Solubilities of the Mo-Si-Ti Phases
Table 3: Invariant Equilibria
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]
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Reaction T [°C] Type Phase Composition (at.%)
Mo Si Ti
395
Landolt-BörnsteinNew Series IV/11A4
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]
Reaction T [°C] Type Phase Composition (at.%)
Mo Si Ti
396
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mo–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Mo
Si Data / Grid: at.%
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
Si Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Mo–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Mo
Si Data / Grid: at.%
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.
Isothermal section at
1600°C [2003Yan]
20
40
60
80
20 40 60 80
20
40
60
80
Ti Mo
Si Data / Grid: at.%
Axes: at.%
β
L
(Ti,Mo)Si2
TiSi2
MoSi2
Mo5Si
3
Mo3Si
TiSi
Ti5Si
4
Ti5Si
3
L
Fig. 7: Mo-Si-Ti.
Isothermal section at
1425°C
[2003Yan, 2004Yan]
401
Landolt-BörnsteinNew Series IV/11A4
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
Si Data / Grid: at.%
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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+MoSi2+SiTiSi2+(Ti,Mo)Si2+Si
(Ti,Mo)Si2+Sie, 1330
e6, 1400
Fig. 9e: Mo-Si-Ti.
Vertical section
through Si
isoconcentrate 60
mass%
404
Landolt-BörnsteinNew Series IV/11A4
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
(Ti,Mo)Si2+MoSi2+SiTiSi2+(Ti,Mo)Si2+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
(Ti,Mo)Si2+MoSi2+SiTiSi2+(Ti,Mo)Si2+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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Temperature – Composition Sections
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.
Notes on Materials Properties and Applications
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
Landolt-BörnsteinNew Series IV/11A4
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,
427-432 (1954) (Crys. Structure, Experimental, 20)
[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.
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)
[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
Landolt-BörnsteinNew Series IV/11A4
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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.
Diagram, Experimental, 6)
[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.
Diagram, Experimental, 27)
[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)
(Experimental, Mechan. Prop., Phys. Prop., 57)
[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)
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Calculated isothermal
section at 1800°C
[1998Zen]
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Landolt-BörnsteinNew Series IV/11A4
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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.
Calculated isothermal
section at 1500°C
[1998Zen]
Fig. 4: N-Ni-Ti.
Calculated isothermal
section at 1353°C
[1998Zen]
413
Landolt-BörnsteinNew Series IV/11A4
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.
Calculated isothermal
section at 1100°C
[1998Zen]
Fig. 6: N-Ni-Ti.
Calculated isothermal
section at 900°C
[1998Zen]
414
Landolt-BörnsteinNew Series IV/11A4
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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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Invariant Equilibria
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].
Temperature – Composition Sections
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.
Notes on Materials Properties and Applications
[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
Landolt-BörnsteinNew Series IV/11A4
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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
(1950) (Experimental, Crys. Structure, 14)
[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,
Crys. Structure, 18)
[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
(1982) (Experimental, Mechan. Prop., 14)
[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)
[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)
[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)
[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)
(Experimental, Equi. Diagram, 11)
[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)
(Crys. Structure, Experimental, 20)
[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
Landolt-BörnsteinNew Series IV/11A4
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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.
Structure, Experimental, 4)
[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)
Table 1: Crystallographic Data of Solid Phases
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
Landolt-BörnsteinNew Series IV/11A4
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N–Ti–V
Table 2: Invariant Equilibria
, 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]
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Calculated isothermal
section at 1800°C and
pN2 = 8.106 bar
[1998Zen]
Fig. 5: N-Ti-V.
Calculated liquidus
surfaces at pN2 = 1bar
[1998Zen]
423
Landolt-BörnsteinNew Series IV/11A4
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.
Calculated isothermal
section at 1200°C and
pN2 = 1 bar [1998Zen]
Fig. 8: N-Ti-V.
Calculated isothermal
section at 1000°C and
pN2 = 8.106 bar
[1998Zen]
424
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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.
Pseudobinary Systems
[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.
Temperature – Composition Sections
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.
Notes on Materials Properties and Applications
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,
Experimental, #, *, 10)
427
Landolt-BörnsteinNew Series IV/11A4
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
(1982) (Crys. Structure, Experimental, 2)
[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),
112-116 (1983) (Crys. Structure, Experimental, 9)
[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),
186-187 (1992) (Equi. Diagram, Experimental, #, *, 4)
[1992Shi] Shimizu, K., Tadaki, T., “Recent Studies on the Precise Crystal-Structural Analyses of
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)
(Crys. Structure, Experimental, 18)
[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
(2000) (Experimental, 13)
428
Landolt-BörnsteinNew Series IV/11A4
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)
Table 1: Crystallographic Data of Solid Phases
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
pure Ti at 25°C [Mas2]
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
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
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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
Pd 50.00Data / Grid: at.%
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.
A partial isothermal
section at 800°C
Fig. 5: Ni-Pd-Ti.
A partial isothermal
section at 700°C
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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
Pd 50.00Data / Grid: at.%
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.
A partial isothermal
section at 400°C
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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
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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
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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.
Pseudobinary Systems
[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].
Invariant Equilibria
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.
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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.
Notes on Materials Properties and Applications
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)
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Landolt-BörnsteinNew Series IV/11A4
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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
(1961) (Experimental, Crys. Structure, 7)
[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,
277-293 (1990) (Review, Equi. Diagram, 11)
[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
Titanium System”, ASM Internal, Metals Park, OH, 197-210 (1987) (Equi. Diagram, Crys.
Structure, Thermodyn., Review, 110)
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Landolt-BörnsteinNew Series IV/11A4
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Ni–Si–Ti
[1987Mur2] Murray, J.L., “The Si-Ti (Silicon-Titanium) System” in “Phase Diagrams of Binary
Titanium System”, ASM Internal, Metals Park, OH, 289-294 (1987) (Equi. Diagram, Crys.
Structure, Thermodyn., Review, 29)
[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)
(Experimental, Crys. Structure, 20)
[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,
Thermodyn., Review, 60)
[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,
Crys. Structure, 112)
[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
(2002) (Experimental, 16)
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Table 1: Crystallographic Data of Solid Phases
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
pure Ti at 25°C [Mas2]
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
Landolt-BörnsteinNew Series IV/11A4
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
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Ni–Si–Ti
Table 2: Invariant Equilibria
* 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]
Reaction T [°C] Type Phase Composition (at.%)
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
Landolt-BörnsteinNew Series IV/11A4
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
Si 50.00Data / Grid: at.%
Axes: at.%
e1
e3
E e2
Ti2Ni
Ti5Si
3
Fig. 2: Ni-Si-Ti.
Partial liquidus
surface
444
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Ni–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Ni
Si Data / Grid: at.%
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
Si Data / Grid: at.%
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.
Isothermal section at
750°C
Fig. 4: Ni-Si-Ti.
Isothermal section at
1000°C
445
Landolt-BörnsteinNew Series IV/11A4
MSIT®
Ni–Si–Ti
20
40
60
80
20 40 60 80
20
40
60
80
Ti Ni
Si Data / Grid: at.%
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.
Isothermal section at
1100°C
top related