Formation of Ti-Zr-Cu-Ni bulk metallic glasses
X. H. Lina) and W. L. Johnson
W M. Keck Laboratory of Engineering Materials, California institute of Technology, Pasadena,
California 91125
{Received 30 May 1995; accepted for publication 17 August 1995)
Formation of bulk metallic glass in quaternary Ti-Zr-Cu-Ni alloys by relatively slow cooling from
the melt is reported. Thick strips of metallic glass were obtained by the method of metal mold
casting. The glass forming ability of the quaternary alloys exceeds that of binary or ternary alloys
containing the same elements due to the complexity of the system. The best glass forming alloys
such as Ti&r,,Cu,,Nis can be cast to at least 4-mm-thick amorphous strips. The critical cooling
rate for glass formation is of the order of 250 K/s or less, at least two orders of magnitude lower than
that of the best ternary alloys. The glass transition, crystallization, and melting behavior of the alloys
were studied by differential scanning calorimetry. The amorphous alloys exhibit a significant
undercooled liquid region between the glass transition and first crystallization event. The glass
forming ability of these alloys, as determined by the critical cooling rate, exceeds what is expected
based on the reduced glass transition temperature. It is also found that the glass forming ability for
alloys of similar reduced glass transition temperature can differ by two orders of magnitude as
defined by critical cooling rates. The origins of the difference in glass forming ability of the alloys
are discussed. It is found that when large composition redistribution accompanies crystallization,
glass formation is enhanced. The excellent glass forming ability of alloys such as Tis,Zr,,Cu,,Nis
is a result of simultaneously minimizing the nucleation rate of the competing crystalline phases. The
temary/quaternary Laves phase (MgZn, type) shows the greatest ease of nucleation and plays a key
role in determining the optimum compositions for glass formation. 0 199.5 American Institute
of
Physics.
I. INTRODUCTION
The reduced glass transition temperature Trs of glass
forming alloys, which is the ratio of glass transition tempera-
ture T8 to melting temperature T,,, , plays an important role in
determining the glass forming ability of alloy~.‘~’ Alloys hav-
ing higher Trp generally exhibit better glass forming ability
(GFA). The critical cooling rate required to avoid the forma-
tion of detectable fraction of crystal in quenching molten
alloys is used in describing-the glass forming ability of ma-
terials. In quenching molten alloys, a sample of typical di-
mension R and initial temperature T, will require a total
cooling time 7 (to T,) of the order of:
r-(R’/K),
(1)
where K is the thermal diffusivity of the alloy. It is given by
K
= KIC where K is the thermal conductivity and C is the
heat capacity per unit volume. The cooling rate achieved will
be of the order of
. dT On,-T,) _ KKn-T,)
T=z= 7 - CR2 .
(2)
Taking T, - T,-400 K, K-O.1 W/cm s-t K-t (typical of a
molten alloys), and C-4 J/cm3 K-t (also typical of molten
alloys), gives
F(K/s)= lOIR’(cm).
(3)
Therefore, the maximum thickness of the amorphous alloy is
determined by the critical cooling rate of the material. Notice
“‘Electronic mail: xhlin@cco.caltech.edu
that one order of magnitude increasing in the maximum di-
mension is equivalent to a two order of magnitude decrease
in the critical cooling rate. Figure 1 shows the critical cool-
ing rates for a number of glass forming alloys as a function
of the reduced glass transition temperature and correspond-
ing maximum dimension. The values of alloys l- 17 in this
figure are taken from Davies.3 The critical cooling rate for
Zr41,2Ti13.sCu12,5Ni10Be22,5 (Ref. 4) is taken from direct
mea-
surement by using an electrostatic levitation facility.* The
critical cooling rates of other alloys are the upper limit val-
ues estimated by taking the value of the maximum thickness
of an amorphous ribbon obtained by melt spinning or half
the value of the reported maximum dimension for glass
phases obtained by normal metal mold casting as maximum
thickness, and by assuming ideal cooling. The actual critical
cooling rates may be lower.
In many alloy systems Tg does not vary with composi-
tion as rapidly as T,,, . Therefore, it is frequently the case that
alloys near deep eutectics have high Trg and, hence, exhibit
better glass forming ability. The Ti-Cu, Zr-Cu, Ti-Ni, and
Zr-Ni systems all exhibit deep eutectics. Binary alloys in
these systems have been known to form glass by rapid so-
lidification for many years. On the other hand, the crystal
structures of the binary intermetallic compounds which form
in these binary alloys differ among the systems. Thus, one
expects to find ternary or quaternary eutectics in- the
Ti-Zr-Cu-Ni system, and the quaternary eutectic melting
temperatures may be lower than those of the binary systems.
As such, alloys near the quatemary eutectics
may
show
greater glass forming ability than corresponding binary al-
loys.
6514 J. Appl. Phys. 78 (il), 1 December 1995
0021-8979/95/78( 11)/6514/6/$6.00
Q 1995 American Institute of Physics
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8
01
0.2 0.3 0.4 0.5 0.6 0.7
Tg/Tm
FIG. 1. Critical cooling rates for glass formation and corresponding maxi-
mum thickness of glass phase. Key to the alloys: (1) Ni; (2) Fes,Bs; (3)
Fe@ll; (4) Te; (5) Au77.&3.$i8.& (6) hdb7; (7) Fe4dJi4dQ7; (8)
Co,sSi,sB,e; (9) Ge; (10) Fe,,Si,eB,,; (11) Ni,sSisB,,; (12) FesaP,sC,; (13)
Pkd%Pz; (14) Pdd%; (15) Ni62,4Nb37:6; (16) Pd77.5Cu6Si16.5; (17)
Pd,eNi,aP, (above from Ref. 3); (18) Aus5Pbs.s5Sbszs (Ref. 6); (19)
h~~~Ni~oCu~0
(Ref. 7); (20) Mg&rssYr, (Ref. 8); (21)
Zr&ut,,Ni,oAl,, (Ref. 9); (22) Zr,,,,Til,,sCu,,,Ni,~e,,,s (Refs. 4 and 5);
~23m~wwi,.
The phase diagrams of ternary Ti-Zr-Ni and Ti-Zr-Cu
have been studied respectively by two different groups. For
the Ti-Zr-Ni system,”
it uras found that for the phase dia-
gram section of Ti,Ni-Zr,Ni, replacing Ti with Zr up to 13.5
at. % in Ti,Ni leads to a lowering of liquidus temperature
from 1273 to 1173 K. Replacing Zr~ with Ti in Zr,Ni up to
22.7 at. % Ti leads to a lowering of liquidus temperature
from 1313 to 1123 K (the melting temperature of Zr2Ni’used
here is different from the 1393 K value given in the binary
ahoy phase diagrams”). From 13.5 to 44 at. % of Zr, the
liquidus temperature rises to 1193 K owing to the appearance
of a new ternary “MgZnz-type” Laves phase. For Ti-Zr-Cu
system,12 near the section of TiCu-ZrCu, three eutectics
were found at Ti34.42zrl7.9aCu47.59,
Ti14.24Zr37.13Cu48.639
and
Tit7.37Zr,3,0Cus,,3. Between these three eutectics, there is
also a ternary MgZn,type Laves phase. Massalski, Woychik,
and Dutkiewicz have reported a very broad glass forming
region for ternary Ti-Zr-Cu alloys by melt spinning.t3 The
thickness of the melt-spun ribbons is about 50 pm, and the
estimated cooling rate is about 5 X 10’ K/s.
We have used a copper mold-casting technique to exam-
ine the glass forming ability of both Ti-Zr-Ni and
Ti-Zr-Cu alloys. We found no Ti-Zr-Ni alloys which can
be cast to the amorphous state using a 300-,um-thick copper
mold. For the Ti-Zr-Cu system, the best glass forming ahoy
was found at Ti,sZr,,Cus,, which can be cast to the amor-
phous state to the thickness of 500 pm. The critical cooling
rate for glass formation is estimated to be 2X lo4 K/s. In the
present article we report two bulk glass forming regions in
the Ti-Zr-Cu-Ni quatemary system. Here, we rather arbi-
trarily define a “bulk” metallic glass as having a minimum
20 30 40 50 60
Atom percent Zr
FIG. 2. Two bulk glass formation regions in the pseudoternary phase dia-
gram. The dots represent the alloys which can be cast to amorphous strips of
at least 1 mm thickness. Generally for the titanium-rich region the nickel
concentration is about 4- 12 at. %. The glass forming region is largest when
nickel concentration is about 8 at %. For the zirconium-rich region, copper
and nickel are roughly interchangeable when at least 4 at. % of either copper
or nickel is used. The region moves downward with increasing nickel con-
centration. In the center of the diagram the quaternary Laves phase field is
shown.
dimension of 1 mm equivalent to a critical cooling rate of
4x103 K/s.
II. EXPERIMENT
Ingots of alloys were prepared by induction melting
99.99% pure Ti, 99.8% pure Zr, 99.999% pure Cu, and
99.97% pure Ni on a water-cooled silver or copper boat un-
der a Ti-gettered argon atmosphere. The nominal composi-
tions are used in the current article. The weight loss of the
samples by alloying was less than 0.1%. Thus, the composi-
tions of the alloys did not change significantly after melting.
The alloy ingots were then remelted under vacuum in a
quartz tube using an rf induction coil and then injected into a
copper mold under pure argon at about 1 atm pressure. The
copper mold has internal strip-shaped cavities of -about 2 cm
length, 4-6 mm width, and varying thickness of 300 pm,
500 pm, 1 mm, 2 mm, 3 mm, and 4 mm. This yields cast
samples of varying strip thickness. The typical length of the
resulting strips is 20 mm. The typical dimensions of the strip
cross sections are 1X4, 2X4, 3X4 or 4X6 mm2. The
crystalline/amorphous nature of the strips was determined by
x-ray diffraction, using a 120” position sensitive detector
(Inel) and a collimated Co KCY x-ray source. To ensure the
amorphous nature of the interior of the strips, some strips
were cut longitudinally in half, and the cross-sectional sur-
faces were examined by x-ray diffraction. The glass transi-
tion and crystallization behavior were studied using a
Perkin-Elmer differential calorimeter (DSC-4). The melting
temperatures of the crystalline alloys were measured using a
Setaram high-temperature differential thermal analysis
(DTA). Vicker’s hardness of the material was obtained using
a Leitz microhardness tester.
Ill. RESULTS
A. Glass forming regions
Figure 2 illustrates two regions of the quaternary com-
position space in which bulk glass formation (>l mm thick-
J. Appl. Phys., Vol. 78, No. 11, 1 December 1995 X. H. Lin and W. L. Johnson 6515
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1~1111l~~~I~~~~Il~~~‘~~~~~~~~~‘~~~~~~~~~’~~
40
60 80
100
120
Two Theta (degree)
FIG. 3. X-ray-diffraction pattern (Co Kcu radiation) taken from the cross-
sectioned surface of 4X6X20 mm3 T$Zr,,Cu,,Nis strip obtained by metal
mold casting.
ness) was found. Notice that the diagram is pseudoternary,
the information regarding the copper/nickel ratio is not seen.
Generally, for the titanium-rich region, the nickel concentra-
tion is about 4-12 at. %. The glass forming region is largest
when the nickel concentration is about 8 at. %. For the
zirconium-rich side, copper and nickel are interchangeable
when one maintains a minimum of 4 at. % of either copper
or nickel. This region tends to shift downward with increas-
ing nickel concentration. The admixture of Cu and Ni im-
proves the GFA tremendously. The best Ti-Zr-Cu amor-,
. .
phous alloy 1s T@r,,Cu,,, having a maximum glass
thickness or about 0.5 mm. Ti,,Zr,,Cu,,Nis is the best qua-
ternary Ti-Zr-Cu-Ni amorphous alloy. It can be cast
af
least
4 mm thick. The estimated critical cooling rate for glass
formation of this alloy is about 250 K/s or lower. It is two
orders of magnitude lower than that of the best ternary
Ti-Zr-Cu glass former. We also found that when an alloy
was cast to a strip thicker than its maximum thickness for
glass formation, the outer layer of the strip remains amor-
phous, while the core is crystalline. This suggests that the
glass forming ability of these alloys is restricted by homoge-
neous nucleation in the sample interior, rather than by het-
erogeneous nucleation at the copper mold interface, if we
assume there are no heterogeneous nucleation sites within
the molten liquid. Figure 3 shows the x-ray-diffraction pat-
tern for one of the strips. No peaks corresponding to crystal-
line phases can be detected.
B. Thermal analysis of the amorphous alloys
Figure 4 shows the DSC traces of two amorphous alloys
taken using a heating rate of 20 Wmin. They exhibit an
endothermic heat event characteristic of the glass transition
followed by four characteristic exothermic heat release
events indicating the successive stepwise transformations
from a metastable undercooled liquid state tci the equilibrium
crystalline intermetallic phases at different temperatures. Tg
is defined as the onset of the glass transition temperature,: TX,
Temperature (K)
FIG. 4.
DSC scans of amorphous alloys. Yfs is the onset of glass transition
temperature, r,, is the onset of first crystallization temperature, and so on.
(4 %~I&USS. Q-4 ‘%4~11CU47Nl
‘8.
is the onset temperature of the first crystallization event, etc.
AT,. defined as TX, -
Tg
, is referred to as the supercooled
liquid region. For the Ti,,Zr,,Cu,,Nis amorphous alloy,
TX=671 K, TX, =717 K, and AT=46 K, respectively. For the
Ti,5Zr&u55, T,=668 K, TX,=697 K, and AT=29 K, re-
spectively. Table I shows glass transition and crystallization
temperatures of
some
of the representative glass forming al-
1oys:Figure 4 shows a high-temperature DTA scan of the
crystalline alloy Ti,~Zr,,Cu,Nis. The alloy begins to melt at
a solidus temperature Tsol=
1105 K followed by complete
melting at the liquidus temperature Tfiq= 1160 K.
C. Mechanical properties
Vicker’s hardness measurements on these amorphous
strips have been carried out. The typical accuracy of the
measurement is 3%. Table II shows the Vicker’s hardness of
some
representative glassy alloys. In terms of hardness, the
composition dependence Ni>Cu@-Zr>Ti is gene&y ob-
served. The value for Ti,,Zr,,Cu,,Nis alloy is H, =628+20
TABLE I. Glass transition and crystallization temperatures of some repre-
sentative glass forming alloys at heating rate of 20 Wmin. The missing
crystallization temperatures are due either to the absence of crystallization
peaks up to 873
K,
or to the crystallization temperature being too close to
873 K to be determined accurately. Here X73 K is the maximum working
temperature of the DSC.
Composition
Ti Zr
Cu Ni T, (K) L, W Tti (K) TX3 (K) TX4 IK)
35 IO
5.5 0
668
697 760 805
33 13.4
49.6 4 674
694 758 800
34 11
47 8 671
717 778 813
33.4 11.9
42.7 12 671
724 780
9.9 43.3
42.8 4 6.57
691
742 791
9.5 45.2
37.3 8 653
690 729 769
9.5 48.9
29.6- 12 637
678 709 748
782
10 50
20 20 644
680 749 759
780
6516 J. Appl. Phys., Vol. 78, No. 11, 1 December 1995
X. H. Lin and W. L. Johnson
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TABLE II. Vicker’s hardness of some representing glass forming alloys
determined using a load ‘of 500 g.
7 I, I I h 1, I I I I, I ! I,,,
Composition
l-l zr
36.9
5.8
33 5.8
33 9.6
34
11
39.6 5.5
35.9 5.5
32.2 5.5
28.5 9.2
32.2 9.2
28.5
12.9
9.5 37.8
9.5 41.5
9.5 45.2
9.5 48.9
10
55
5.8 45.2
5.8 48.9
5 55
10
50
10
50
CU Ni
53.3 4-
57.2 4 i
53.4 ~- 4
47 8
46.9 8
._
50.6 8
54.3 8
54.3 8
50.6 8
SO.6 8
44.7 8
41 8
37.3 8
33.6 8
25 10
11 8
37.3 8
32 8
20 20
1.5 25
H,, (kg/md
562
626
618
628
596
616
642
670
618
626
616
553
519
501
462
536
523
487
499
558
, I
I I9 I I I I I I I I r 9 8 I I
1050
1100 1150 1100
Temperature (K)
FIG. 5. High-temperature DTA scan of Ti&ZrL,Cu471$g. T,, is the solidus
temperature. Tti, is the liquidus temperature.
kg/mm’. Using the well-known’relation H, 7 3 a,, l4 , the
yield strength of this amorphous alloy is estimated to be
around 2 GPa.
One 2-mm-thick Ti,,Zr,,Cu,Nis amorphous strip was
successively rolled at room temperature using a thickness
reduction of 1.5% deformation per step down to a 0.15~mm-
thick ribbon without cracking. This demonstrates the ductile
behavior of the amorphous material when deformation oc-
curs under a confined geometry. The resulting ribbon can be
further bent 180” without failure. The Vicker’s hardness of
the resulting ribbon was also measured and agrees with that
of the initial strip within the experiment accuracy. This indi-
cates that there is no work hardening as is-expected for an
amorphous material.
IV. DISCUSSION
In Fig. 1, we see a strong Trg dependence of the glass
forming ability of metallic alloys. High values of Trs- are
associated with good glass forming ability. For example,
Pd4aNi4aP2a alloy has Trg=0.66, s,15 Aus5Pb22,5Sb22 s alloys
have Tra=0.63.6 Their critical cooling rates are 10’ K/s or
less. The recently found Zr41,2Ti13.8Cu,15Ni,0Be225 alloys
have Trg=0.67,4 and have critical cooling rate of 1 K/s5 The
rrs of Ti,4Zr,1Cu47Ni8 is 0.578, by comparison not a very
high value. According to Fig. 1, the estimated critical cooling
rate for an alloy of T,=O.578 would typically be of the order
of lo5 K/s. Therefore, from the point of view of reduced
glass transition temperature, the GFA of Ti-Cu-Zr-Ni al-
loys such as Ti34ZrllCu47Nis is .relatively better than ex-
pected (see Fig. 5). In fact, this is not unique. Some newly
found multicomponent alloys, such as Mgs&!uz,Y,, (Ref. 8)
and Zr,&ui7.5Ni10A1~J (Ref. 9) also show better GFA than
expected from Trg. Three factors, a significantly different
atomic size among the constituent elements, a large negative
heat of mixing, and the necessity of substantial redistribution
of the component elements for the progress of crystallization,
have been cited and used to interpret the unexpected good
GFA of these systems.i6 In fact, the first two factors, large
atomic size ratios and large negative heat of mixing, are
already reflected by the relatively low lying eutectic melting
temperatures of these alloys. As such; these factors give, at
best, a qualitative glass forming criteria, whose role is not
quantified. In conventional homogeneous nucleation theory,
one evaluates GFA by considering the competition between a
decreasing- nucleation barrier and an increasing viscosity
with increasing undercooling of the liquid. If the nucleation
crystalline phase has a composition very different from that
of the undercooled liquid, only when the composition of a
local liquid region the size of a critical crystalline nucleus
satisfies the composition requirements of the crystalline
phase (either by fluctuation of liquid decomposition) can
crystallization occur. For higher-order multicomponent sys-
tems, it is more difficult for the concentrations of all ele-
ments to simultaneously satisfy the composition require-
ments of crystalline phase than for lower-order systems. As
such, the crystallization process of the multicomponent un-
dercooled liquid will tend to be more sluggish than for sim-
pler systems. The multicomponent alloys may thus exhibit
better GFA than predicted from the point of view of the
reduced glass transition temperature alone. The fundamental
argument used here has been loosely called the “confusion
principle.” I7 The concept is that as the number of compo-
nents in a liquid alloy is increased, crystallization becomes
confused or frustrated. Recently, Desre has quantified this
argument.18 He considered a multicomponent system with II
equally concentrated component elements. For n.= 2 to
y1= 10, the additions of each additional component to the al-
loy is predicted to lower by an order of magnitude the prob-
ability of the concentration fluctuation within a cluster of 600
atoms for crystallization. The probability of achieving a criti-
cal nucleus of the required composition is lowered by an
order of magnitude with the addition of each new compo-
J. Appl. Phys., Vol. 78, No. 11, 1 December 1995
X. H. Lin and W. L. Johnson 6517
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nent. Thus, the more complex the alloy, the better the GFA.
This argument may explain the extraordinary good GFA of
the present Ti-Zr-Cu-Ni system and the Zr-Cu-Ni-Al
(Ref. 9) and Mg-Cu-Y (Ref. 8)
systems as well.
One piece
of evidence in support of this argument is that all good glass
forming alloys in these systems show a large undercooled
liquid region, as indicated by the temperature interval be-
tween glass transition and first crystallization event during
sample heating. The nucleation of crystalline phases from the
undercooled liquid phases requires substantial atomic diffu-
sion and composition redistribution. In this situation, the un-
dercooled liquid is relative stable against the crystallization.
Note that for the Ti34Zr11Cu47Nis amorphous alloy AT=46
K, while for the Ti,,Zr,,Cu,, alloy AT=29 K. AT correlates
with the relative GFA of these two alloys.
It is noteworthy that a large number of early transition-
metal/late-transition-metal systems exhibit the MgZn, or the
MgCu,-type Laves phases. Examples are: TiMn?, ZrMn,,
TiF%, and TiZn, (of MgZnz-type structure), and ZrFez,
ZrCo,, and ZrZn, (of MgCu,-type structure). The congruent
melting temperatures of these Laves phases are generally
above 1600 K and the liquidus curves are relatively flat (with
respect to composition). The Laves phase homogeneity
ranges are often as large as 10 at. % or more.‘i These factors
suggest that the Gibbs free-energy function of the Laves
phases varies slowly with composition. In this sense, the
Laves phases are somewhat similar to liquid or glassy phase.
They are both “forgiving” of composition variation. On the
other hand, there are no equilibrium binary Ti-Cu, Ti-Ni,
Zr-Cu, or Zr-Ni Laves phases. The melting temperature of
Ti,Ni is lowered by adding Zr, and the melting temperature
of Zr,Ni is lowered by adding Ti. One would expect an ex-
tremely low lying liquidus temperature in the center of the
Ti2Ni-Zr2Ni section if there were no ternary Laves phase. In
the actual ternary Ti-Ni-Zr
system, the
Laves phase enters
the diagram with a homogeneity range from 21 to 30 at. %
Zr at 970 K.” One expects similarly large homogeneity
ranges for the ternary Ti-Cu-Zr Laves. phase as well as the
quaternary Ti-Zr-Cu-Ni Laves phase. Because of the rela-
tive insensitivity of the Laves phase to composition varia-
tions, it
may
nucleate more easily than the other competing
crystalline phases from.the undercooled liquid. We. have de-
termined that there are two quaternary eutectics near the
compositions- of Ti,,Zr,,Cu,Ni, and Ti,,Zr,Cu,sNi,, , re-
spectively. Interestingly, these two alloys cannot be cast to
0.5~mm-thick amorphous
strips. By comparison,
Ti,&r,,Cu47Ni8 can be cast to 4-mm-thick amorphous strips
and Ti&r4,Cu,sNi8 can be cast to at least 2-mm-thick amor-
phous strips. -Their reduced glass transition temperatures are
in fact lower than those of the two eutectic alloys, but their
critical cooling rates for glass formation are at least one and
possibly two orders of magnitude lower than those of the two
eutectic alloys. They are much better glass formers. Wesug-
gest that this is due to the co.mpositions being relatively far
from that of the Laves phase. To obtain a good glass forming
composition in Ti-Zr-Cu-Ni system, one must consider
two factors: First, one should be close to the eutectic points
to obtain a high reduced glass transition temperature; second,
one must avoid the Laves phase which apparently nucleates
relatively. more easily than other competing crystalline
phases. The best glass forming alloys of the present Ti-Zr-
Cu-Ni system are examplessof simultaneously satisfying
these two conditions.
The main obstacle which prevents us from getting better
GFA in the Ti-Zr-Cu-Ni system is the existence of the
quaternary Laves phase. To improve the GFA of the
Ti-Zr-Cu-Ni
system,
one must find a way to eliminate or at
least destabilize the Laves phase to achieve a lower melting
temperature alloy in the center of the quasiternary phase dia-
gram. We have found that transition metals having a high
melting temperature Laves phase with Ti or Zr tend to stabi-
lize the Laves phase of the quatemary alloy and thus degrade
the GFA. On the other hand, substituting Cu or Ni by Zn or
Co does not degrade the glass forming ability. If the fifth
element added is metalloid element such as B or Si, the alloy
loses its GFA. In this case, crystallization is apparently trig-
gered by the precipitation of very stable Ti or Zr borides or
silicides. It seems the only way to improve the glass forming
ability of current Ti-Zr-Cu-Ni system is to add Be as has
already been proven.4,5 Apparently the very small atomic ra-
dius of Be is incompatible with the preferred atomic size
ratio of the Laves phase. As such, Be acts to destabilize the
quaternary Laves phase resulting in further depression of the
alloy liquidus curve in the center portion of the quatemary
alloy diagram. This yields a pentiary alloy with a eutectic at
943 K2. From this point of view, we can understand why the
Zr-Ti-Cu-Ni-Be system is such an exceptional glass
former.
From the above arguments, one concludes that while the
reduced glass transition temperature still plays a dominant
role in determining the GFA of metallic alloys, the require-
ment that the crystallization involve large composition Buc-
tuations in the liquid phase also tends to enhance the GFA
tremendously.
V. CONCLUSIONS
Bulk metallic glass formation in the ternary
Ti-Zr-Cu-Ni system is reported. Bulk samples of metallic
glass can be prepared by metal mold casting up to dimen-
sions of several millimeters. The critical cooling rate for
glass formation is of the order of 500 K/s. For the particular
amorphous alloy Ti34ZrllCU47Ni8, the hardness is about 628
kg/mm2, while the tensile strength is estimated to be about 2
GPa. Comparing with the Zr-Ti-Cu-Ni-Be system, the.
quaternary alloys have relatively poorer GFA. On the other
hand, the absence of Be in these glasses may make them of
interest from the point of view of applications.
It has been noted that the GFA of the quaternary alloys is
better than might be expected based on their reduced glass
transition temperature alone. This suggests that the require-
ment for composition fluctuations and diffusion controlled
nucleation play an important role in determining the GFA of
alloys.
The greatly enhanced GFA obtained by adding a few
percent of a fourth element in the present alloy system sug-
gests that making a system more complex is a practical way
to search for new bulk glass forming alloy systems.
6518 J. Appt. Phys., Vol. 78, No. 11, 1 December 1995
X. H. Lin and W. L. Johnson
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