metals
Article
Formation and Corrosion Behavior of
Mechanically-Alloyed Cu–Zr–Ti Bulk
Metallic Glasses
Pee-Yew Lee 1 , Yeh-Ming Cheng 1 , Jyun-Yu Chen 1 and Chia-Jung Hu 2, *
1
2
*
Institute of Materials Engineering, National Taiwan Ocean University, Keelung 20224, Taiwan;
[email protected] (P.-Y.L.); [email protected] (Y.-M.C.); [email protected] (J.-Y.C.)
Department of Materials Engineering, Tatung University, 40 Chungshan N. Road, 3rd Sec.,
Taipei 10452, Taiwan
Correspondence: [email protected]; Tel.: +886-2-2182-2928 (ext. 6228); Fax: +886-2-2593-6897
Academic Editors: Jordi Sort Viñas and K.C. Chan
Received: 31 August 2016; Accepted: 18 April 2017; Published: 20 April 2017
Abstract: Cu60 Zr30 Ti10 metallic glass powder was prepared by mechanically alloying a mixture
of pure Cu, Zr, and Ti powders after 5 h of milling. Cu60 Zr30 Ti10 bulk metallic glass (BMG) was
synthesized by vacuum hot pressing the as-milled Cu60 Zr30 Ti10 metallic glass powder at 746 K in
the pressure range of 0.72–1.20 GPa, and the structure was analyzed through X-ray diffraction and
transmission electron microscopy. The pressure could enhance the thermal stability, and prolong
the existence, of the amorphous phase inside the Cu60 Zr30 Ti10 powder. Furthermore, the corrosion
behavior of the Cu-based BMG in four corrosive media was studied using a potentiodynamic method.
The Cu60 Zr30 Ti10 BMG exhibited a low corrosion rate and current density in 1 N solutions of H2 SO4 ,
NaOH, and HNO3 . X-ray photoelectron spectroscopy results revealed that the formation of Zrand Ti-rich passive oxide layers provides a high corrosion resistance against 1 N H2 SO4 and HNO3
solutions, and the breakdown of the protective film by Cl− attack was responsible for pitting corrosion
in a 3 wt % NaCl solution. The formation of oxide films and the nucleation and growth of pitting
were analyzed through microstructural investigations.
Keywords: mechanical alloying; bulk metallic glass; corrosion; supercooled liquid region; vacuum
hot pressing
1. Introduction
Recently, new metallic glass alloys with a wide supercooled liquid region exceeding 30 K have
been prepared in numerous Cu-based alloy systems, such as Cu–Ti–Zr, Cu–Hf–Ti, Cu–Zr–Hf–Ti,
Cu–Zr–Ti–Y, Cu–Ti–Zr–Ni–Si, and Cu–Ti–Zr–Ni–Sn systems [1–3]. Cu-based bulk metallic glasses
(BMGs) with diameters ranging from 2 mm to 6 mm have been produced using conventional casting
techniques at a low cooling rate [2]. These new alloys are expected to increase the number of
application fields of BMGs because of their unique properties, such as high tensile strength, large
elastic limits (~2%), and high resistance to corrosion and wear. The defining characteristic of metallic
glasses, compared to conventional metallic materials, is their lack of crystallinity, and associated
lack of microstructural features, such as grain and phase boundaries [4]. Metallic glasses clearly
cannot have the crystallographically-defined slip-systems of polycrystalline metals. In the absence
of dislocation-mediated slip, they show high yield stresses, much closer to the theoretical limit than
their crystalline counterparts [5]. Therefore, the lack of grain structure and associated microstructural
features (e.g., solute segregation) gives corrosion resistance. BMGs exhibited excellent corrosion
resistance compared to that of the crystalline alloys. This can be attributed to the beneficial effect of the
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alloying components, such as Nb in ZrCuNiAl BMG [6], or Ag, Ni in passive film on the MgCu BMG [7],
which are homogeneously mixed in the amorphous phase. Thus far, many researchers have studied
the glass-forming ability, thermal stability, and mechanical properties of all of the aforementioned
Cu-based BMGs. While BMGs are still little used because of their macroscopic brittle nature and the
difficulty of processing, thin film metallic glasses (TFMGs) are a possible solution to take advantage of
the unique MG properties of high strength [8,9], large plasticity within ∆T, excellent wear resistance,
annealing-induced amorphization, and the resulting smooth surfaces [10]. The lack of ductility is
mainly due to the detrimental formation of a few large localized shear bands, leading to the catastrophic
fracture [11,12]. There have been several approaches to improve the room temperature ductility in
BMGs, including microstructure modification by forming composite structures in the amorphous
matrix [13,14], and surface modifications, such as thick (~90 µm) coatings [15]. However, data on the
corrosion behavior of Cu-based BMGs and TFMGs are comparatively rare. The application of Cu-based
BMGs requires that they have an acceptable lifetime in different environments. Hence, investigating the
corrosion behavior of Cu-based metallic glasses in many aggressive environments, such as acidic and
salt solutions, at ambient temperatures and in humid air, is of scientific and technological consequence.
A large supercooled liquid region is known to be essential for the formation of BMGs through
rapid solidification; however, this restricts BMG formation to near-eutectic compositions, where
supercooling can be achieved without the nucleation of crystalline phases. In this regard, an alternative
method to prepare metallic glasses is mechanical alloying (MA) [16]. Previous investigations have
demonstrated the successful use of MA for synthesizing metallic glass powders with large supercooled
regions [17,18]. The existence of a wide supercooled liquid region enables the preparation of BMGs by
hot pressing or extruding metallic glass powders at temperatures above the glass transition temperature
(Tg ) because of the rapidly decreasing viscosity of the metallic glass [19,20]. Therefore, in this study, the
feasibility of preparing a Cu60 Zr30 Ti10 BMG by using the MA and vacuum hot pressing techniques was
investigated. Furthermore, the corrosion behavior of the Cu60 Zr30 Ti10 BMG in four corrosive media
(1 N H2 SO4 , HNO3 , and NaOH and 3 wt % NaCl) was studied using a potentiodynamic method.
2. Experimental Procedure
Elemental powders of Cu (99.9%, <300 mesh), Ti (99.9%, <100 mesh), and Zr (99.9%, <325 mesh),
were weighed to yield the desired composition, Cu60 Zr30 Ti10 , and then canned into an SKH 9
high-speed steel vial together with Cr steel balls in an argon-filled glove box, where a SPEX 8016 shaker
(SPEX SamplePrep, Metuchen, NJ, USA) ball mill was used for MA. The details of the MA process
are described elsewhere [21]. The bulk samples were prepared by consolidating the as-milled powder
in a vacuum hot pressing machine (Giant-asia, Miaoli, Taiwan) at 746 K under a pressure range of
0.72–1.20 GPa. The structure of the as-milled and bulk samples was analyzed using a Siemens D-5000
X-ray diffractometer (CountrySiemens AG, Karlsruhe, Germany). Thermal analysis was performed
using a Dupont 2000 differential scanning calorimeter (DSC, CountryTA instruments, New Castle,
PA, USA) at a heating rate of 40 K/min. The corrosion characteristics of the alloys were determined
on the basis of anodic polarization curves obtained using a potentiodynamic method. Prior to the
corrosion tests, the specimens were mechanically polished in cyclohexane with silicon carbide paper
up to 2000 grit, degreased in acetone, washed in distilled water, dried in air, and further exposed
to air for 24 h to ensure high reproducibility. The potentiodynamic polarization test was performed
using the three-electrode principle with a Solartron model 1286A potentiostat (TICS International Ltd.,
Hounslow, UK), where a platinum electrode was used as the counter electrode and a saturated calomel
electrode as the reference electrode. The samples were tested in 1 N solutions of H2 SO4 , HNO3 , and
NaOH and a 3 wt % NaCl solution at room temperature. The anodic polarization curves were measured
at a potential sweep rate of 1 mV/s after open-circuit immersion for approximately 30 min when the
open-circuit potential became almost steady. The morphology of the corroded surface was examined
by scanning electron microscopy (SEM, Hitachi High-Technologies Corporation, Tokyo, Japan) and
energy dispersive X-ray microanalysis (Hitachi High-Technologies Corporation, Tokyo, Japan) as
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performed using a Hitachi S-4800. X-ray photoelectron spectroscopy (XPS, VG Scientific , Waltham,
performed
using a Hitachi
X-ray
photoelectron
spectroscopy(XPS,
(XPS, VGScientific
Scientific
, Waltham,
performed
a HitachiS-4800.
S-4800.
X-ray
photoelectron
spectroscopy
, Waltham,
MA, USA) using
measurements
for
surface
analysis
of the specimens
before and VG
after the electrochemical
MA,
USA)
measurements
for
surface
analysis
of
the
specimens
before
and
after
the
electrochemical
test
MA,
USA)
forsolutions
surface analysis
of the specimens
and after
the electrochemical
test in
the measurements
aforementioned
were performed
using abefore
VG ESCA
Scientific
Theta Probe
in the
aforementioned
solutionssolutions
were performed
using a VG
ESCA
Theta
ProbeTheta
photoelectron
test
in the aforementioned
were performed
a Scientific
VG(hESCA
Scientific
Probe
photoelectron
spectrometer with
monochromatized
Al Kausing
excitation
ν = 1486.6 eV).
spectrometer
with
monochromatized
Al Ka excitation
1486.6 eV).
photoelectron
spectrometer
with monochromatized
Al(hKa
(hν = 1486.6 eV).
ν =excitation
3. Results and Discussion
3. Results
and
Discussion
3. Results
and
Discussion
3.1. Ball-Milled Powder and Consolidated BMG
3.1.3.1.
Ball-Milled
Powder
Ball-Milled
Powderand
andConsolidated
ConsolidatedBMG
BMG
Figure 1 depicts the X-ray diffraction pattern of the Cu60Zr30Ti10 powder milled for 5 h. A broad
Figure
1 depicts
the X-ray
diffraction
pattern
ofofthe
Cu
Ti1010powder
powdermilled
milledfor
for
5 h.
broad
30
Figure
1 depicts
X-ray
diffraction
pattern
the
Cu6060Zr
Zr
30Ti
5 h.
AA
broad
diffraction
peak
at 2θthe
= 40°
is the
main feature
of this
pattern,
indicating
that the as-milled
powder
is
◦ is the main feature of this pattern, indicating that the as-milled powder is
diffraction
peak
at at
2θ2θ
= 40
diffraction
peak
=
40°
is
the
main
feature
of
pattern,
indicating
that
the
as-milled
powder
is
amorphous. The thermal stability of the amorphous Cu60Zr30Ti10 powder was investigated through
amorphous.
thermal
stability
Cu
Ti1010powder
powder
wasinvestigated
investigated
through
amorphous.
The
thermal
stabilityofofthe
theamorphous
amorphous
Cu6060Zr30
30Ti
was
differentialThe
scanning
calorimetry,
and
the
corresponding
scans
are depicted
in Figurethrough
2.
The
differential
scanning
calorimetry,
and
the
corresponding
scans
are
depicted
in
Figure
2.
The
amorphous
differential
scanning
calorimetry,
and
the
corresponding
scans
are
depicted
in
Figure
2.
The
amorphous powder exhibited an endothermic heat event because of the glass transition; this event is
amorphous
exhibited
anheat
endothermic
heat event
of
the glass
transition;
event
isa a
powder
exhibited
endothermic
eventindicating
because
of successive
thebecause
glass transition;
this
event is this
followed
followed
bypowder
a an
sharp
exothermic
peak,
stepwise
transformation
from by
followed
by
a
sharp
exothermic
peak,
indicating
successive
stepwise
transformation
from
a
sharp
exothermic
peak,
indicating
successive
stepwise
transformation
from
a
supercooled
liquid
state
supercooled liquid state to crystalline phases. The glass transition temperature (Tg) and the
supercooled
liquid
state
to
crystalline
phases.
The
glass
transition
temperature
(T
g
)
and
the
to crystalline
phases.
The glass(Ttransition
temperature
(Tg ) and the The
crystallization
(TΔT
x ) are
crystallization
temperature
x) are 736 K
and 779 K, respectively.
supercooledtemperature
liquid region
x
temperature
(T
x) are
736
K and 779
K, Cu
respectively.
supercooled
liquid
ΔT
x
736crystallization
779 K,is
respectively.
The
supercooled
liquid
region
∆T10xBMG
isThe
43 (i.e.,
K, which
to K
the[23]).
rapid
isK43and
K, which
similar to the
rapid
solidification
of
60Zr30Ti
ΔTx = is
42similar
K [22],region
45
is
K,ΔT
which
is
to
the
rapid
solidification
60Zr30Ti10 BMG (i.e., ΔTx = 42 K [22], 45 K [23]).
solidification
ofknown
Cusimilar
Ti10
BMG
(i.e.,
42of
K Cu
[22],
45 K(GFA)
[23]). [24],
As the
∆Tbe
known
as one
As43the
x is
as
of
indices
of∆T
glass
ability
it can
said
the alloy
has of
x =forming
x is
60 Zr
30one
As
the
ΔT
x is known as one of indices of glass forming ability (GFA) [24], it can be said the alloy has
very of
high
glass
forming
ability.
indices
glass
forming
ability
(GFA) [24], it can be said the alloy has very high glass forming ability.
very high glass forming ability.
Cu60Zr30Ti10/MA 5h
Cu60Zr30Ti10/MA 5h
Figure 1. X-ray diffraction pattern of mechanically-alloyed Cu60Zr30Ti10 powder after 5 h of milling.
Figure
1. X-ray
diffraction
pattern
ofofmechanically-alloyed
Zr30TiTi1010powder
powder
after
h of
milling.
Figure
1. X-ray
diffraction
pattern
mechanically-alloyedCu
Cu60
60Zr30
after
5 h5 of
milling.
Cu60Zr30Ti10/MA 5h
Cu60Zr30Ti10/MA 5h
Figure 2. DSC trace of amorphous Cu60Zr30Ti10 powder after 5 h of milling.
Figure
2. DSC
trace
amorphousCu
Cu60Zr
Zr30Ti10 powder after 5 h of milling.
Figure
2. DSC
trace
ofofamorphous
60 30 Ti10 powder after 5 h of milling.
The powder milled for 5 h was consolidated into a disk by using a vacuum hot press operated at
The
powder
milled pressures
for 5 h was consolidated
into
a disk
using 3a depicts
vacuumahot
pressconsolidated
operated at
746
K
30 min
under
0.72, 0.96, and
1.20
GPa.by
Figure
typical
The for
powder
milled
for 5 h wasof
consolidated
into
a disk
by
using a vacuum
hot press
operated at
746
K
for
30
min
under
pressures
of
0.72,
0.96,
and
1.20
GPa.
Figure
3
depicts
a
typical
consolidated
sample
ofmin
BMG
that
exhibited
a smooth
outer
surface
and 3metallic
luster.
Itsconsolidated
polished,
746sample
K for 30of
under
pressures
of
0.72,
0.96,
and
1.20
GPa.
Figure
depicts
a
typical
BMG
that
exhibited through
a smooth
outer
surface inand
metallic
luster.pores
Its nor
polished,
cross-sectional
view,
as examined
SEM,
is depicted
Figure
4. Neither
voids
sample
of BMG that
exhibited
a smooth
outerSEM,
surface
and metallic
luster.4.ItsNeither
polished,
cross-sectional
cross-sectional
view,
as examined
through
is depicted
in Figure
pores
nor voids
view, as examined through SEM, is depicted in Figure 4. Neither pores nor voids were observed, which
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were observed, which indicates that bulk Cu60Zr30Ti10 with a highly dense structure was successfully
were observed, which indicates that bulk Cu60Zr30Ti10 with a highly dense structure was successfully
indicates
thatTransmission
bulk Cu60 Zr30electron
Ti10 with
a highly dense
structure
was successfully
prepared.status
Transmission
prepared.
microscopy
was used
to examine
the amorphization
of the
prepared. Transmission electron microscopy was used to examine the amorphization status of the
electron
microscopy
was
used
to
examine
the
amorphization
status
of
the
corresponding
disk
corresponding bulk disk sample. The HRTEM bright field image is shown in Figurebulk
5. No
corresponding bulk disk sample. The HRTEM bright field image is shown in Figure 5. No
sample.
The HRTEM
field
is shown
in Figure
5. No
appreciable
contrast
revealing
appreciable
contrast bright
revealing
theimage
presence
of crystalline
phases
is observed
in the
bulk sample,
appreciable contrast revealing the presence of crystalline phases is observed in the bulk sample,
theindicating
presenceitofis crystalline
phases
is observed
in The
the bulk
it is a homogeneous
a homogeneous
amorphous
phase.
insertsample,
shown inindicating
Figure 5 presents
the selected
indicating it is a homogeneous amorphous phase. The insert shown in Figure 5 presents the selected
area diffraction
pattern
recorded
from
the 5matrix,
shows area
a typical
amorphous
amorphous
phase. The
insert
shown in
Figure
presentswhich
the selected
diffraction
patternpattern
recorded
area diffraction pattern recorded from the matrix, which shows a typical amorphous pattern
characterized
a diffuse
ring. amorphous
The corresponding
diffraction
these
from
the matrix,by
which
showshalo
a typical
pattern X-ray
characterized
by patterns
a diffuseof
ring.
characterized
by
a diffuse
halo
ring. The corresponding
X-ray
diffraction
patterns
ofhalo
these
hot-pressed
samples
arediffraction
depicted inpatterns
Figure 6.
All
of the
samples exhibit
amorphous-representing
The
corresponding
X-ray
of
these
hot-pressed
samples
are
depicted
in
Figure
hot-pressed samples are depicted in Figure 6. All of the samples exhibit amorphous-representing 6.
broad
peaks.amorphous-representing
However, on a decrease
in the
pressurepeaks.
to 0 GPa
(samples
only in
Allbroad
of the diffraction
samples exhibit
broad
diffraction
However,
on awith
decrease
diffraction
peaks. However, on a decrease
in the
pressure to 0 GPa
(samples
with
only
sintering
treatment),
the
X-ray
diffraction
pattern
exhibits
several
crystalline
peaks
at
about
40–45°of
thesintering
pressuretreatment),
to 0 GPa (samples
with
only
sintering
treatment),
the
X-ray
diffraction
pattern
exhibits
the X-ray diffraction pattern exhibits several crystalline peaks at about 40–45°of
2θ with
the broad
peak
representing
the2θamorphous
phase.peak
This result indicates
that the
several
crystalline
peaks
at about
40–45◦ of
with the broad
the amorphous
2θ with
the broad
peak
representing
the amorphous
phase. Thisrepresenting
result indicates
that the
consolidation
pressure
suppressed
crystallization
in the Cu
60Zr30Ti10 BMG. The crystallization
phase.
This result
indicates
that thethe
consolidation
pressure
suppressed
the crystallization
in the
consolidation
pressure
suppressed
the
crystallization
in the Cu
60Zr30Ti10 BMG.
The crystallization
behaviors of BMG under high pressure have been extensively studied [25,26]. In general, the
Cubehaviors
Zr
Ti
BMG.
The
crystallization
behaviors
of
BMG
under
high
pressure
have
been
extensively
of
BMG
under
high
pressure
have
been
extensively
studied
[25,26].
In
general,
the
60 30 10
application pressure on BMG has three possible effects [27]. The first effect is densification, which
application
pressure
on BMG
has three possible
effects
[27]. has
The three
first effect
is densification,
which
studied
[25,26].
In general,
the application
pressure
on BMG
possible
effects [27]. The
first
favors the appearance of the closest structure and the crystallization process. The second effect is
favors
the appearance
of thefavors
closestthe
structure
and the
crystallization
process.and
Thethe
second
effect is
effect
is densification,
which
appearance
of the
closest structure
crystallization
suppression of atomic mobility, which reduces the atomic diffusion in metallic glasses. The third
suppression
of atomic
mobility,
which reduces
the mobility,
atomic diffusion
in metallic
The third in
process.
The second
effect
is suppression
of atomic
which reduces
the glasses.
atomic diffusion
effect is due to changes in the relative Gibbs free energies of the glassy phase and crystalline phases.
effect glasses.
is due toThe
changes
the relative
free in
energies
of theGibbs
glassyfree
phase
and crystalline
phases.
metallic
thirdin
effect
is due toGibbs
changes
the relative
energies
of the glassy
phase
In the present study, increasing the pressure to ≥0.72 GPa hindered long-range atomic diffusion in
In
the
present
study,
increasing
the
pressure
to
≥0.72
GPa
hindered
long-range
atomic
diffusion
in
and
crystalline
phases.
In
the
present
study,
increasing
the
pressure
to
≥
0.72
GPa
hindered
long-range
the BMG because the diffusion barrier was elevated by the reduction of volume and the
the
BMG
because
the
diffusion
barrier
was
elevated
by
the
reduction
of
volume
and
the
atomic
diffusion
the BMG
becausewas
the caused
diffusion
was elevated
by thethat
reduction
of volume
and
annihilation
ofin
excess
free volume
bybarrier
high pressure.
This means
the second
effect of
annihilation
of
excess
free
volume
was
caused
by
high
pressure.
This
means
that
the
second
effect
of
thesuppression
annihilationofofatomic
excessmobility
free volume
was pressures
caused bybecomes
high pressure.
This means
that
thethe
second
effect
at higher
a dominating
factor;
thus,
amount
suppression of atomic mobility at higher pressures becomes a dominating factor; thus, the amount
of suppression
of
atomic
mobility
at
higher
pressures
becomes
a
dominating
factor;
thus,
the
amount
of the amorphous phase in the hot-pressed samples increases with increasing pressure. Similar
of the amorphous phase in the hot-pressed samples increases with increasing pressure. Similar
of the
amorphous
phase
in theinhot-pressed
with
pressure.
results
have been
observed
BMGs such samples
as Al84Niincreases
10Ce6, Al82.5
Ni5Yincreasing
8Co2Zr2.5 [28],
and Zr55Similar
Al10Ni5results
Cu30
results have been observed in BMGs such as Al84Ni10Ce6, Al82.5Ni5Y8Co2Zr2.5 [28], and Zr55Al10Ni5Cu30
have
been observed in BMGs such as Al84 Ni10 Ce6 , Al82.5 Ni5 Y8 Co2 Zr2.5 [28], and Zr55 Al10 Ni5 Cu30 [29].
[29].
[29].
Figure
Sampleofofaahot-pressed
hot-pressed Cu60Zr
Ti
BMG
disk.
Figure
3. 3.
Sample
Zr30
Ti10
BMG
disk.
Figure
3.
Sample of a hot-pressedCu
Cu60
60Zr
30
1010
BMG
disk.
30Ti
Figure
4. Polished
cross-sectionalview
view of Cu6060Zr
Zr30Ti
746
KK
under
a 1.20
GPa
Figure
4. Polished
cross-sectional
Ti1010BMG
BMGconsolidated
consolidatedat
746
under
a 1.20
GPa
Figure
4. Polished
cross-sectional viewofofCu
Cu60Zr30
30Ti10
BMG
consolidated
atat746
K under
a 1.20
GPa
pressure.
pressure.
pressure.
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Figure 5. TEM bright-field image and selected area diffraction pattern of Cu60Zr30Ti10 BMG
Figure 5. TEM bright-field image and selected area diffraction pattern of Cu60 Zr30 Ti10 BMG consolidated
consolidated
746 K underimage
a pressure
1.20 GPa.
Figure
5. TEMatbright-field
and ofselected
area diffraction pattern of Cu60Zr30Ti10 BMG
at 746 K under a pressure of 1.20 GPa.
Relative Intensity (arb. units)
Relative Intensity (arb. units)
consolidated at 746 K under a pressure of 1.20 GPa.
Cu60Zr30Ti10Bulk 746K 30min
Cu60Zr30Ti10Bulk 746K 30min
1.2 GPa
1.2 GPa
0.96 GPa
0.96 GPa
0.72 GPa
0.72 GPa
0 GPa
0 GPa
20
20
30
30
40
40
50
50
2θ
2θ
60
60
70
70
80
80
Figure 6. X-ray diffraction patterns of Cu60Zr30Ti10 BMG consolidated at 746 K under different
pressures.
Figure
6. X-ray
diffraction
patterns
of 60Cu
Ti10
BMG
consolidated
K under
different
Figure
6. X-ray
diffraction
patterns
of Cu
Zr6030ZrTi3010
BMG
consolidated
at at
746746
K under
different
pressures.
pressures.
3.2. Corrosion Behavior of Cu–Zr–Ti BMG Alloys
3.2. Corrosion Behavior of Cu–Zr–Ti BMG Alloys
The polarization curves for Cu60Zr30Ti10 BMG consolidated at different pressures are depicted
The
polarization
for
Zr TiTi
BMG
consolidated
different
pressures
are depicted
60
10
in The
Figures
7. Some curves
crucial
including
corrosion
potential
(Eat
c), corrosion
current
(ic),
polarization
curvesdata
for Cu
Cu
60Zr3030
10the
BMG
consolidated
at different
pressures
are density
depicted
in Figure
7. potential
Some
crucial
data
including
thedensity
corrosion
potential
), corrosion
density
cthe
(Ep), and
passive
current
(ip)potential
derived
from
curves
arecurrent
summarized
in passive
Figures
7. Some
crucial
data
including
the
corrosion
(E
c),(E
corrosion
current
density
(ic), in (ic ),
passive
potential
current
density
from
the
Table
1. From (E
the
curves
in Figure
7, (i
inp)(iacid
(1 N from
H2SO
4 and
1 Ncurves
HNO
3)are
andsummarized
alkaline
passive
potential
(E
ppolarization
),and
and passive
passive
current
density
derived
the
curves
are summarized
in (1 in
p ),
p ) derived
N1.NaOH)
solutions,
all of the
curves
are
S-type
aN
of
potential
versus
the
Table
1. From
thepolarization
polarization
curves
ininFigure
7, in
acid
(1 on
N(1H
2plot
SOH4 2and
1
N
HNO
3
)
and
alkaline
(1
Table
From
the
curves
Figure
7,curves
in acid
SOelectrode
and
1
N
HNO
)
and
alkaline
4
3
logarithm
of
the
current
density.
Most
of
the
test
samples
show
an
obvious
passive
region
in
anodic
N
NaOH)
solutions,
all
of
the
curves
are
S-type
curves
on
a
plot
of
electrode
potential
versus
the
(1 N NaOH) solutions, all of the curves are S-type curves on a plot of electrode potential versus the
−5 A·cm
−2. Since
polarization
and
the ipdensity.
is
between
10−3of
and
10
noan
pitting
corrosion
was
observed
in
logarithm
the
current
density.
Most
the
test
show
obvious
passive
region
in anodic
logarithm
of of
the
current
Most
of
the
testsamples
samples
show
an obvious
passive
region
inthe
anodic
−3 and
−5 A·cm
−2to
alloys polarized
inipacid
solution10at
a
potential
up
0.5
V,
the
alloys
are
suggested
to
tend
to
form
polarization
and the
is between
10
.
Since
no
pitting
corrosion
was
observed
in
the
polarization and the ip is between 10−3 and 10−5 A·cm−2 . Since no pitting corrosion was observed
stable
protective
films immediately
1N
2SO
1 N HNO
solutions.to
However,
for the
alloys
polarized
in surface
acid solution
at a potentialinup
toH
0.5
V,4 and
the alloys
are 3suggested
tend to form
in the alloys polarized in acid solution at a potential up to 0.5 V, the alloys are suggested to tend to
samples
in 1 Nsurface
NaOHfilms
solutions,
at a potential
of2SO
approximately
0.5 3V,
the anodic
currentfor
density
stable
protective
immediately
in 1 N H
4 and 1 N HNO
solutions.
However,
the
form
stable in
protective
surface
films which
immediately
inapproximately
1 N H2ofSO
1 N HNO
However, for
4 and
3 solutions.
increases
withsolutions,
potential,at
could be
transpassive
dissolution
of the alloys.
samples
1rapidly
N NaOH
a potential
of because
0.5 V, the
anodic
current
densityAt
theincreases
samples
in 1 Nofwith
NaOH
solutions,
at could
a potential
of scattered,
approximately
0.5
V, the
current
the beginning
polarization,
the
current
values
are
as represented
byanodic
two
or alloys.
three
active
rapidly
potential,
which
be because
of transpassive
dissolution
of the
Atdensity
increases
rapidly
with
potential,
which
could
be
because
of
transpassive
dissolution
of
the
alloys.
peaks.
Metastable pitting
corrosion
onare
surface
defects,
such as micropores
during
thecurrent
beginning
of polarization,
the current
values
scattered,
as represented
by two orformed
three active
At the
beginning
of
polarization,
the
current
values
are
scattered,
as
represented
by
two
or
three
consolidation
or the competition
between on
thesurface
formation
of active
and passive
ZrO2 and
TiOactive
2
current
peaks. Metastable
pitting corrosion
defects,
suchCuO
as micropores
formed
during
consolidation
or the competition
the on
formation
active CuO
ZrO2 formed
and TiO2during
current
peaks. Metastable
pittingbetween
corrosion
surfaceofdefects,
suchand
as passive
micropores
3.2. Corrosion Behavior of Cu–Zr–Ti BMG Alloys
consolidation or the competition between the formation of active CuO and passive ZrO2 and TiO2
surface layers, may have caused this scattering. A detailed study on the appearance of this specific
Metals 2017, 7, 148
Metals 2017, 7, 148
6 of 12
6 of 12
behavior
is underway.
In the
3 wtthis
% NaCl
solution,
the samples
pressures
surface
layers, may have
caused
scattering.
A detailed
study on consolidated
the appearanceatoflower
this specific
(0.72 behavior
and 0.96 is
GPa),
underwent
dissolution
immediately
after
the open circuit
Ecpressures
was reached,
underway.
In theactive
3 wt %
NaCl solution,
the samples
consolidated
at lower
− -containing
(0.72because
and 0.96of
GPa),
active dissolution
immediately
after
open circuitsolutions.
Ec was reached,
possibly
the underwent
unstable surface
films of the
alloys in the
Clthe
However,
−
possiblyconsolidated
because of theat
unstable
surface films
the alloys
the Cl
-containing
solutions.
However, with
in a sample
high pressure
(1.20ofGPa),
the in
alloy
was
spontaneously
passivated
−2pressure
sample consolidated
at·cm
high
(1.20ofGPa),
the alloy was
spontaneously
passivatedresistance
with
an ip in
of aapproximately
10−3 A
. The effect
consolidation
pressure
on the corrosion
an ip of approximately 10−3 A·cm−2. The effect of consolidation pressure on the corrosion resistance of
of Cu60 Zr30 Ti10 BMG was evaluated by comparing the Ec and ic listed in Table 1. An increase in
Cu60Zr30Ti10 BMG was evaluated by comparing the Ec and ic listed in Table 1. An increase in pressure
pressure was observed to reduce the ic . This implies that the Cu60 Zr30 Ti10 BMG consolidated with
was observed to reduce the ic. This implies that the Cu60Zr30Ti10 BMG consolidated with high
high pressure
exhibitsimproved
improved
corrosion
resistance
against
various
corrosive
solutions.
Theare
results
pressure exhibits
corrosion
resistance
against
various
corrosive
solutions.
The results
are asasexpected
because
the
density
of
Cu
Zr
Ti
BMG
increases
with
pressure
during
consolidation.
60
30
10
expected because the density of Cu60Zr30Ti10 BMG increases with pressure during consolidation.
Furthermore,
compared
with
thethe
casting
samples
Zr3030TiTi
BMG
[22],
passive
region
Furthermore,
compared
with
casting
samplesof
ofCu
Cu60
60Zr
1010
BMG
[22],
thethe
passive
region
of theof the
potentiodynamic
polarization
curves
undulation
N2SO
H24SO
HNO
potentiodynamic
polarization
curvesexhibits
exhibits an
an unsettled
unsettled undulation
in in
thethe
1 N1 H
and
HNO
3
4 and
3
solutions.
In 3the
wtNaCl
% NaCl
solution,aapitting
pitting attack
attack was
InIn
acid
(1 N
and4 1and
N 1N
solutions.
In the
wt3 %
solution,
wasnot
notobserved.
observed.
acid
(1H
N2SO
H24 SO
3) and alkaline (1 N NaOH) solutions, the surface morphology of the Cu60Zr30Ti10 BMG
HNOHNO
3 ) and alkaline (1 N NaOH) solutions, the surface morphology of the Cu60 Zr30 Ti10 BMG exhibited
exhibited
passive state
according
to thecurve,
polarization
curve, representing
theofoccurrence
of pit
a passive state aaccording
to the
polarization
representing
the occurrence
pit corrosion
on the
corrosion on the sample surface. These results suggest that the passive films formed on the casting
sample surface. These results suggest that the passive films formed on the casting sample surface are
sample surface are more unstable in alkaline solutions than in 1 N H2SO4 and HNO3, and 3 wt %
more unstable in alkaline solutions than in 1 N H2 SO4 and HNO3 , and 3 wt % NaCl solutions after
NaCl solutions after polarization tests, indicating that a weaker surface film is responsible for the
polarization
tests, indicating
lower corrosion
resistance.that a weaker surface film is responsible for the lower corrosion resistance.
Figure 7. Anodic and cathodic polarization curves of Cu60Zr30Ti10 BMG consolidated at 746 K under
Figure 7. Anodic and cathodic polarization curves of Cu60 Zr30 Ti10 BMG consolidated at 746 K under
different pressures in (a) 1 N H2SO4, (b) 1 N NaOH, (c) 1 N HNO3, and (d) 3 wt % NaCl solutions
different pressures in (a) 1 N H2 SO4 ; (b) 1 N NaOH; (c) 1 N HNO3 , and (d) 3 wt % NaCl solutions open
open to air at 298 K.
to air at 298 K.
Table 1. Polarization parameters of Cu60Zr30Ti10 BMG consolidated at 746 K under different
H2SO
4, 1 N NaOH,
3 wt % NaCl,
and
1N
solutions
open to air at 298
SEMpressures
imagesinof1 Nthe
corroded
surfaces
of Cu
Ti10 3BMG
consolidated
at K.
746 K under a
60 Zr
30HNO
1.2 GPa pressureItem
in different solutions open to air at 298
K
are
depicted
in
Figure
8. The SEM
Pressure (GPa)
Ec (Volts)
ic (A/cm2)
Ep (Volts)
ip (A/cm2)
Solution
observations were recorded after potentiodynamic polarization to determine whether a pitting attack
0.72
−0.05
8.4 × 10−6
>0.5 V
2.2 × 10−4
had occurred.
In acid (1 N0.96
H2 SO4 and −0.05
1 N HNO3 ) and
alkaline (1>0.5
NV
NaOH) solutions,
the surface
1 N H2SO4
6.9 × 10−6
6.4 × 10−5
−6
−5 the glassy
morphology of the Cu60 Zr1.2
Ti
BMG
alloy
did
not
exhibit
obvious
pits,
indicating
that
>0.5
V
2.2
×
10
−0.07
1.2
×
10
30 10
alloy, after polarization in these solutions, appeared to have a homogeneous surface morphology.
Moreover, from the polarization curves of the glassy alloy in these solutions, the Cu60 Zr30 Ti10 BMG
alloy appeared highly resistant to corrosion, or the corrosion rate decreased because of the formation
Metals 2017, 7, 148
7 of 12
of uniform passive films. By contrast, the surface morphology of the Cu60 Zr30 Ti10 BMG alloy after the
polarization
test
in the 3 wt % NaCl solution was very different from that after polarization in7 the
Metals 2017,
7, 148
of 12other
−
three solutions, indicating that the surface is not uniformly corroded in the Cl -containing solution.
0.72
4.8 × 10−4
>0.5 V
10−3
The surface morphology was
not even for−0.55
the BMG alloy
after polarization
in the 3 wt1.8%× NaCl
solution;
1 N NaOH
0.96
−0.56
2.2 × 10−5
>0.5 V
6.5 × 10−5
−5
−5
many deep corrosion pits were
dispersed
at
the
surface
(Figure
8d).
Figure
9
depicts
SEM
micrographs
<0.5 V
3.1 × 10
1.2
−0.52
1.6 × 10
0.72 Cu Zr Ti0.27BMG alloy
4.9consolidated
× 10−6
--of the corroded surface of the
at---746 K under different
pressures
60 30 10
3 wt % NaCl
0.96
−0.37
4.3 × 10−6
----in 3 wt % NaCl solutions. In
all
of
the
cases,
characteristics
of
the
active
nature,
represented
by deep
<0.5 V
1.1 × 10−3
1.2
−0.34
1.5 × 10−6
−4
−4
corrosion pits, were observed
the
and size
0.72 in these glassy
0.04 alloys in2.8
× 10NaCl solution.
>0.5 V Moreover, the
5.7 ×range
10
1 N HNO3
−5
0.96 increased 0.03
7.9 × 10
V 1.2 to 0.721.5GPa.
× 10−4 This result
of corrosion pitting gradually
with a decrease
in pressure>0.5
from
indicates that the Cu60 Zr30 Ti10 BMG alloy was corroded in the NaCl solution because of the presence
SEM images of the corroded surfaces of Cu60Zr30Ti10 BMG consolidated at 746 K under a 1.2 GPa
of Cl− ; rapid, extremely-localized pit growth occurred, leading to the breakdown of the passive film,
pressure in different solutions open to air at 298 K are depicted in Figure 8. The SEM observations
as evidenced
by theafter
pitspotentiodynamic
formed in the NaCl
solution
9). whether
As indicated
in Li’s
study
were recorded
polarization
to (Figure
determine
a pitting
attack
had [30],
BMGsoccurred.
containing
high
Cu
content
usually
exhibit
poor
corrosion
resistance
in
chloride
ion-containing
In acid (1 N H2SO4 and 1 N HNO3) and alkaline (1 N NaOH) solutions, the surface
solutions
due to the
formation
of10soluble
porous
Cu–Cl
films.
The dissolution
of elemental
Cu led to
morphology
of the
Cu60Zr30Ti
BMG alloy
did not
exhibit
obvious
pits, indicating
that the glassy
alloy,
after polarization
in these solutions,
appeared
to have a homogeneous surface morphology.
pitting
corrosion
and the breakdown
of the passive
film.
Moreover, from the polarization curves of the glassy alloy in these solutions, the Cu60Zr30Ti10 BMG
Table
1. Polarization
parameters
Cu60 Zr30 Ti
BMG
consolidated
at 746 K under
different
pressures
alloy appeared
highly
resistant toofcorrosion,
or10the
corrosion
rate decreased
because
of the formation
in
N H2 SOpassive
,
1
N
NaOH,
3
wt
%
NaCl,
and
1
N
HNO
solutions
open
to
air
at
298
K.
of 1uniform
films.
By
contrast,
the
surface
morphology
of
the
Cu
60
Zr
30
Ti
10
BMG
alloy
after the
4
3
polarization test in the 3 wt % NaCl solution was very different from that after polarization in the
Item
other three solutions,
indicating that the surface is not uniformly corroded in the Cl−-containing
Ep (Volts)
ip (A/cm2 )
Ec (Volts)
Pressure (GPa)
ic (A/cm2 )
solution.
The
surface
morphology
was not even
for the BMG
alloy after polarization in
the 3 wt %
Solution
NaCl solution; many deep corrosion
pits
were
dispersed
at
the
surface
(Figure
8d).
Figure
9 depicts
0.72
−0.05
>0.5 V
8.4 × 10−6
2.2 × 10−4
−
6
5
SEM micrographs
of
the
corroded
surface
of
the
Cu
60
Zr
30
Ti
10
BMG
alloy
consolidated
at
746
K −under
1 N H2 SO4
0.96
−0.05
>0.5 V
6.9 × 10
6.4 × 10
−
6
−
5
1.2
−
0.07
>0.5
V
1.2
×
10
2.2
×
10
different pressures in 3 wt % NaCl solutions. In all of the cases, characteristics of the active nature,
−
4
−
represented by deep corrosion0.72
pits, were observed
glassy alloys
−0.55 in these
>0.5in
V the NaCl
4.8 × 10
1.8 × solution.
10 3
−
5
0.96
−
0.56
>0.5
V
2.2 ×
10
6.5in
× pressure
10−5
1 Nthe
NaOH
Moreover,
range and size of corrosion pitting gradually
increased
with a decrease
−5
5
1.2
−
0.52
<0.5
V
1.6
×
10
3.1
×
10−NaCl
from 1.2 to 0.72 GPa. This result indicates that the Cu60Zr30Ti10 BMG alloy was corroded in the
0.72
— occurred, leading
—
4.9 × 10−6 pit growth
solution because of the presence
of Cl−; rapid, 0.27
extremely-localized
to
0.96
−
0.37
— solution (Figure
—
4.3formed
× 10−6 in the NaCl
3
wt
%
NaCl
the breakdown of the passive film, as evidenced by the pits
9).
1.2
−0.34
<0.5 V
1.5 × 10−6
1.1 × 10−3
As indicated in Li’s study [30], BMGs containing high Cu
content usually exhibit poor
corrosion
−
4
4
0.72 solutions due
0.04 to the formation
>0.5 Vporous5.7
2.8 × 10
× 10−films.
resistance
inHNO
chloride
ion-containing
of
soluble
Cu–Cl
1N
3
−
5
−
0.96
0.03
>0.5 V
7.9 × 10
1.5 × 10 4
The dissolution of elemental Cu led to pitting corrosion and
the breakdown of the passive
film.
Figure 8. SEM images depicting the characteristic morphology of Cu60Zr30Ti10 BMG consolidated at
Figure 8. SEM images depicting the characteristic morphology of Cu60 Zr30 Ti10 BMG consolidated at
746 K under a 1.2 GPa pressure and subsequently corroded in different solutions open to air at 298 K
746 K under a 1.2 GPa pressure and subsequently corroded in different solutions open to air at 298 K
(a) 1 N H2SO4, (b) 1 N NaOH, (c) 1 N HNO3, and (d) 3 mass% NaCl.
(a) 1 N H2 SO4 , (b) 1 N NaOH, (c) 1 N HNO3 , and (d) 3 mass% NaCl.
Metals 2017, 7, 148
8 of 12
Metals 2017, 7, 148
8 of 12
Metals 2017, 7, 148
(a)
8 of 12
(b)
(c)
Figure 9. SEM images of the corroded surfaces of Cu60Zr30Ti10 BMG consolidated at 746 K under (a)
(b)Cu60 Zr30 Ti10 BMG consolidated (c)
Figure 9. SEM(a)
images of the corroded surfaces of
at 746 K under
1.2 GPa, (b) 0.96 GPa, and (c) 0.72 GPa pressures in 3 wt % NaCl solutions open to air at 298 K.
(a) 1.2
GPa,
(b)
0.96
GPa,
and
(c)
0.72
GPa
pressures
in
3
wt
%
NaCl
solutions
open
to
air
at 298(a)
K.
Figure 9. SEM images of the corroded surfaces of Cu60Zr30Ti10 BMG consolidated at 746 K under
1.2an
GPa,
(b) 0.96 GPa,
and (c) 0.72 GPa
in 3 wt % NaCl
solutions
open to air at 298
For
improved
understanding
of pressures
the surface-related
chemical
characteristics
of K.
the alloy, the
For
an
improved
understanding
of
the
surface-related
chemical
characteristics
the alloy,
specimens were analyzed through XPS after polarization tests in 1 N H2SO4, 1 N HNO3, 1 of
N NaOH,
For
an
improved
understanding
of
the
surface-related
chemical
characteristics
of
the
alloy,
the
theand
specimens
were analyzed
afterofpolarization
tests in 1 alloys
N H2 SO
NaOH,
3 wt % NaCl
solutions. through
The XPSXPS
spectra
the glassy Cu–Zr–Ti
exhibited
identical
4 , 1 N HNO
3 , 1 Npeaks
specimens were analyzed through XPS after polarization tests in 1 N H2SO4, 1 N HNO3, 1 N NaOH,
and
wt %ofNaCl
solutions.
The
XPS
spectra of tests
the glassy
Cu–Zr–Ti
exhibited
peaks
to3those
Cu, Zr,
Ti, and O
after
polarization
in these
solutionsalloys
(Figure
10). In allidentical
of the cases,
and 3 wt % NaCl solutions. The XPS spectra of the glassy Cu–Zr–Ti alloys exhibited identical peaks
copperofpeaks
were
for the
Cu-containing
Thesolutions
Zr 3d spectrum
setscases,
of
to those
Cu, Zr,
Ti,detected
and O after
polarization
testsalloy.
in these
(Figureconsists
10). In of
alltwo
of the
to those of Cu, Zr, Ti, and O after polarization tests in these solutions (Figure 10). In all of the cases,
3d5/2 and
3d3/2
peaks
representing
theCu-containing
oxidized state.alloy.
The peaks
at 182.05–182.60
and
copper
peaks
were
detected
for the
The located
Zr 3d spectrum
consists
of184.43–
two sets
copper peaks were detected for the Cu-containing
alloy. The Zr 3d spectrum consists of two sets of
4+ ions. Two peaks of Ti 2p1/2 and 2p3/2 from the Ti4+
eV
were
assigned
to
Zr
3d
electrons
of
Zr
of 185.03
3d3d
and
3d
peaks
representing
the
oxidized
state.
The
peaks
located
at
182.05–182.60
5/2
3/2
5/2 and 3d3/2
peaks representing the oxidized state. The peaks located at 182.05–182.60 and 184.43–and
state
appear
at
464.06–464.84
eV and
458.46–458.94
eV.
TheTwo
O 1s spectrum
consists
peaks
184.43–185.03
eV were
assigned
3d electrons
Zr4+
ions.
ofand
Ti 2p
and of
2p3/2
from
185.03 eV were
assigned
to Zrto
3dZr
electrons
of Zr4+ofions.
Two
peaks ofpeaks
Ti 2p1/2
2p1/2
3/2 from the
Ti4+
4+
originating
from
oxygen
in
metal–O–metal
bonds,
metal–OH
bonds,
or
bound
water.
Relatively
the Ti
state
appear
at
464.06–464.84
eV
and
458.46–458.94
eV.
The
O
1s
spectrum
consists
of
peaks
state appear at 464.06–464.84 eV and 458.46–458.94 eV. The O 1s spectrum consists of peaks
small N 1s
andoxygen
Cl 2p in
spectra
were observed
formetal–OH
the specimens
inor
1 bound
N HNO
3 and 3 wt % NaCl
originating
from
metal–O–metal
bonds,
bonds,
water.
Relatively
small
originating
from oxygen
in metal–O–metal
bonds,
metal–OH
bonds,
or bound
water.
Relatively
3−
−
solutions,
respectively,
and
they
were
assigned
to
NO
and
Cl
ions
on
the
surface
film,
respectively.
N 1s
Cl 2p
spectra
were observed
for the specimens
in 1 N
HNO
and%3NaCl
wt %solutions,
NaCl
N 1ssmall
and Cl
2p and
spectra
were
observed
for the specimens
in 1 N HNO
33 wt
3 and
The
peaks of
Cu 2p, Zr and
3d, and
2p assigned
correspond
to
species
theon
oxide
state on
therespectively.
surface film.
3− and
− 3−
−in
solutions,
respectively,
they Ti
were
toCl
NO
andon
Clthe
ions
thefilm,
surface
film,
respectively,
and
they were assigned
to NO
ions
surface
respectively.
The peaks
Similar
peaksofhave
been
in2p
Ti–Zr–Pd–Cu–Sn
[31] and
[32]
The
Cu 2p,
Zrobserved
3d, and Ti
to species
inNi–Nb–Ti–Zr
theon
oxide
onBMGs.
the surface
film.
of Cu
2p,peaks
Zr 3d,
and
Ti 2p
correspond
tocorrespond
species in the
oxide state
the state
surface
film.
Similar
peaks
Similar
peaks
have
been
observed
in
Ti–Zr–Pd–Cu–Sn
[31]
and
Ni–Nb–Ti–Zr
[32]
BMGs.
have been observed in Ti–Zr–Pd–Cu–Sn [31] and Ni–Nb–Ti–Zr [32] BMGs.
Figure 10. Cont.
Metals 2017, 7, 148
Metals 2017, 7, 148
9 of 12
9 of 12
Figure
10. XPS
analysis
of the
surface
filmsfor
forthe
theCu-based
Cu-based glassy
glassy alloy
2SO
4, 14 ,N
Figure
10. XPS
analysis
of the
surface
films
alloyinin1 1NNHH
1 NaOH,
N NaOH,
2 SO
3, and
wtNaCl
% NaCl
solutions
after
polarization
tests,
Cu2p,
2p,(b)
(b)Zr
Zr3d,
3d,(c)
(c)Ti
Ti2p,
2p, and
and (d) O 1s.
1 N HNO
1 N HNO
3 wt3%
solutions
after
polarization
tests,
(a)(a)
Cu
3 , and
1s.
Figure
11 depicts the cationic contents of the surface films for the Cu-based glassy alloy in
Figure 11 depicts the cationic contents of the surface films for the Cu-based glassy alloy in 1 N
1NH
HNO
1 NaOH,
N NaOH,
% NaCl
solutions
after
the polarization
After
3 ,N
H22SO
SO44,,11NNHNO
3, 1
andand
3 wt 3%wt
NaCl
solutions
after the
polarization
test. Aftertest.
the tests
in the
tests all
in of
allthese
of these
solutions,
the
content
of
Zr
in
the
surface
films
of
the
glassy
alloy
was
higher
than
solutions, the content of Zr in the surface films of the glassy alloy was higher than that of
that of
other
elements
in
the
glassy
alloy.
The
content
of
Cu
in
the
surface
film
was
slightly
lower
other elements in the glassy alloy. The content of Cu in the surface film was slightly lower than that
than of
that
Ti in
the(H
acid
HNO3 ) solutions,
whereas
content
Cu in film
the surface
Ti of
in the
acid
2SO4(H
and
HNO
3) solutions,
whereas the
contentthe
of Cu
in theofsurface
was
2 SO
4 and
than that
ofthat
Ti inof
theTiNaOH
NaCland
solutions.
observations
indicate that the
passivethat
film higher
was higher
than
in theand
NaOH
NaCl These
solutions.
These observations
indicate
film is composed
of ZrO2 and
TiO22 in
theTiO
H22SO
HNO
and
ZrO2 and and
CuOZrO
in the
the passive
film is composed
of ZrO
and
in4 and
the H
and HNO
2 SO34 solutions
3 solutions
2 and
NaOH
and
NaCl
solutions.
Thus,
the
Cu
60
Zr
30
Ti
10
BMG
alloy
exhibits
different
corrosion
CuO in the NaOH and NaCl solutions. Thus, the Cu60 Zr30 Ti10 BMG alloy exhibits different corrosion
mechanisms
in different
solutions.
Furthermore,the
theXPS
XPS spectra
spectra reveal
passive
films
formed
mechanisms
in different
solutions.
Furthermore,
revealthat
thatthe
the
passive
films
formed
after electrochemical polarization are primarily enriched in ZrO2 and TiO2 with a small amount Cu
after electrochemical polarization are primarily enriched in ZrO2 and TiO2 with a small amount Cu
oxides for the Cu60Zr30Ti10 BMG alloy. This observation can be explained by the preferential
oxides for the Cu60 Zr30 Ti10 BMG alloy. This observation can be explained by the preferential oxidation
oxidation of Zr and Ti, because elemental Zr and Ti have a high negative heat of oxide formation
of Zr and Ti, because elemental Zr and Ti have a high negative heat of oxide formation compared with
compared with elemental Cu [31].
elementalZhang
Cu [31].
[33] asserted that the high negative heat of formation is a synergetic effect of Ti and Zr
Zhang
[33] asserted
high negative
heat of formation
is a synergetic
effect of Ti
that inhibits
corrosionthat
andthe
improves
the anti-electrochemical
degradation
characteristics
forand
all Zr
that inhibits
corrosion
and
improves
the
anti-electrochemical
degradation
characteristics
for
all
alloys
alloys containing them. The fact that Zr and Ti oxides are chemically more stable and structurally
containing
The
factcan
that
Zr andforTithe
oxides
are chemically
more stable
and
structurally
denser them.
than Cu
oxide
account
enhanced
corrosion resistance
of the
BMG
alloy [34].denser
By
enhancement
related tocorrosion
the difference
in the of
ionizing
energy
of [34].
the outer-shell
than contrast,
Cu oxidethe
can
account formay
the be
enhanced
resistance
the BMG
alloy
By contrast,
electrons in themay
atoms
constituent
Theinlower
ionizing energy
electrons
the enhancement
beof
related
to theelements.
difference
the ionizing
energyofofthe
theouter-shell
outer-shell
electrons
makes
them
more
chemically
active
than
other
electrons
of
the
atom.
Zr
and
Ti
atoms
have
a lower
in the atoms of constituent elements. The lower ionizing energy of the outer-shell electrons
makes
energy (6.84
and 6.82
respectively)
of Cu
[34]. This
themionizing
more chemically
active
thaneV,
other
electrons than
of thethat
atom.
Zr(7.72
and eV)
Ti atoms
havefacilitates
a lower faster
ionizing
formation
of 6.82
a passive
film in the presence
Ti,eV)
and[34].
subsequently
improves
corrosion
energy
(6.84 and
eV, respectively)
than thatofofZr
Cuand
(7.72
This facilitates
faster
formation
resistance and reduces the corrosion rate in this BMG system.
of a passive film in the presence of Zr and Ti, and subsequently improves corrosion resistance and
However, Figure 11 illustrates that Ti and Zr cations are further concentrated in the surface
reduces the corrosion rate in this BMG system.
film. By contrast, the content of Cu cations obviously decreases in the surface after the polarization
However, Figure 11 illustrates that Ti and Zr cations are further concentrated in the surface film.
By contrast, the content of Cu cations obviously decreases in the surface after the polarization test
in all solutions. Where are the Cu ions of the sample surface in these solutions after electrochemical
Metals 2017, 7, 148
10 of 12
test in all solutions. Where are the Cu ions of the sample surface in these solutions after
10 of 12
electrochemical polarization tests? They may be present in the corrosion environment; fast
dissolution of the active component Cu leads to an accumulation of Zr and Ti cations rapidly at the
alloy-solution
interface
initial
of the electrochemical
polarization
test;
therefore,
polarization
tests?
They during
may be the
present
in stage
the corrosion
environment; fast
dissolution
of the
active
spontaneousCu
passivation
The passive
highly
enriched
in alloy-solution
Zr and Ti cations
and
component
leads to anoccurs.
accumulation
of Zr films
and Tiare
cations
rapidly
at the
interface
largely
deficient
in
Cu
cations
after
the
passive
films
are
formed.
This
prevents
further
dissolution
of
during the initial stage of the electrochemical polarization test; therefore, spontaneous passivation
Cu through
the surface
[31].
occurs.
The passive
films
are highly enriched in Zr and Ti cations and largely deficient in Cu cations
the Cu
60Zr
30Ti
10 BMG This
alloyprevents
did not undergo
stable passivation
in the the
NaOH
and[31].
NaCl
after In
thefact,
passive
films
are
formed.
further dissolution
of Cu through
surface
−-containing solution. Figure 9 depicts deep corrosion pits, which are
solutions,
particularly
in
the
Cl
In fact, the Cu60 Zr30 Ti10 BMG alloy did not undergo stable passivation in the NaOH and NaCl
characteristic
of the active
of the Cusolution.
60Zr30Ti10 Figure
BMG alloy
in the
NaCl
solution
after
the
solutions,
particularly
in thenature
Cl− -containing
9 depicts
deep
corrosion
pits,
which
polarization
test.
The
development
of
these
pits
can
be
explained
on
the
basis
of
galvanic
corrosion
are characteristic of the active nature of the Cu60 Zr30 Ti10 BMG alloy in the NaCl solution after the
because of the
combination
of a noble
element
as explained
Cu (+0.35 on
V) and
a base
as Zr
polarization
test.
The development
of these
pitssuch
can be
the basis
of element
galvanicsuch
corrosion
(−1.53 V)of
orthe
Ti (−1.63
V) for Cu–Zr
or Cu–Ti
alloys.
Due
the
compositional
heterogeneities
between
because
combination
of a noble
element
such
asto
Cu
(+0.35
V) and a base
element such
as Zr
Cu
and
Zr
or
Cu
and
Ti
in
the
glassy
matrix,
which
facilitate
the
nucleation
of
pits
[35],
Cu-rich
(−1.53 V) or Ti (−1.63 V) for Cu–Zr or Cu–Ti alloys. Due to the compositional heterogeneities between
regions
the
metal-film
after pitwhich
formation
in Cu
60Zr
30Ti10. Enhanced
Cu
and form
Zr oratCu
and
Ti in theinterface
glassy matrix,
facilitate
the
nucleation
of pitslocal
[35], selective
Cu-rich
dissolution
metals (Zrinterface
and Ti) occurs
some weak
or
chemical
defects,
regions
formofatactive
the metal-film
after pitatformation
in sites
Cu60with
Zr30 Tiphysical
.
Enhanced
local
selective
10
− ions. Continuous corrosion of local sites can also lead
as
a
result
of
an
autocatalytic
reaction
with
Cl
dissolution of active metals (Zr and Ti) occurs at some weak sites with physical or chemical defects, as
the dissolution
of elemental
Cu, which
is−re-deposited
promptly,
forming
a porous
Further
atoresult
of an autocatalytic
reaction
with Cl
ions. Continuous
corrosion
of local
sitesmass.
can also
lead
corrosion
attack
and
pit-depth
growth
proceed
predominantly
through
the
mesh
mainly
driven
by
to the dissolution of elemental Cu, which is re-deposited promptly, forming a porous mass. Further
the
dissolution
of
chloride-active
Zr
and
Ti
[36].
These
phenomena
can
explain
why
the
Cu
60
Zr
30
Ti
10
corrosion attack and pit-depth growth proceed predominantly through the mesh mainly driven by the
BMG alloy of
has
higher corrosion
resistance
in H2SO
4 and HNO
3 solutions
than
NaOH
and NaCl
dissolution
chloride-active
Zr and
Ti [36]. These
phenomena
can
explain why
theinCu
60 Zr30 Ti10 BMG
solutions.
alloy has higher corrosion resistance in H2 SO4 and HNO3 solutions than in NaOH and NaCl solutions.
Metals 2017, 7, 148
Figure 11. Cationic contents in the surface films for the Cu-based glassy alloy in 1 N H2SO4, 1 N
Figure 11. Cationic contents in the surface films for the Cu-based glassy alloy in 1 N H2 SO4 , 1 N HNO3 ,
HNO3, 1 N NaOH, and 3 wt % NaCl solutions after polarization tests.
1 N NaOH, and 3 wt % NaCl solutions after polarization tests.
4. Conclusions
4. Conclusions
Cu60Zr30Ti10 metallic glass powder was prepared successfully through MA after 5 h of milling.
Cu60 Zr30 Ti10 metallic glass powder was prepared successfully through MA after 5 h of milling.
For the Cu60Zr30Ti10 metallic glass powder, the glass transition and crystallization onset
For the Cu60 Zr30 Ti10 metallic glass powder, the glass transition and crystallization onset temperatures
temperatures were 736 K and 779 K, respectively. The supercooled liquid region temperature range
were 736 K and 779 K, respectively. The supercooled liquid region temperature range (∆T) was 43 K.
(ΔT) was 43 K. The as-milled Cu60Zr30Ti10 powder of the end product was consolidated into a disk
The as-milled Cu60 Zr30 Ti10 powder of the end product was consolidated into a disk with a diameter
with a diameter and thickness of 10 and 3 mm, respectively, at 746 K for 30 min under a pressure
and thickness of 10 and 3 mm, respectively, at 746 K for 30 min under a pressure range of 0.72–1.20 GPa,
range of 0.72–1.20 GPa, using the vacuum hot pressing technique. The consolidation pressure
using the vacuum hot pressing technique. The consolidation pressure suppressed crystallization in the
suppressed crystallization in the Cu60Zr30Ti10 BMG. Furthermore, the electrochemical behavior and
Cu60 Zr30 Ti10 BMG. Furthermore, the electrochemical behavior and surface composition analysis of the
surface composition analysis of the Cu60Zr30Ti10 metallic glassy alloy were investigated. The
Cu60 Zr30 Ti10 metallic glassy alloy were investigated. The Cu60 Zr30 Ti10 BMG consolidated using high
Cu60Zr30Ti10 BMG consolidated using high pressure exhibited high corrosion resistance in various
pressure exhibited high corrosion resistance in various corrosion solutions. This result is as expected
because the density of Cu60 Zr30 Ti10 BMG increases with pressure during consolidation. In addition,
Metals 2017, 7, 148
11 of 12
the Cu-based BMG exhibited excellent corrosion resistance after electrochemical tests in the H2 SO4
and HNO3 solutions. The glassy alloy is passivated with a significantly low current density of the
order of 10−4 –10−5 A·cm−2 and a wide passive potential region in the H2 SO4 and HNO3 solutions,
indicating high corrosion resistance of the glassy alloy in these solutions. The high corrosion resistance
of the alloy is attributed to its chemically homogeneous nature and the passive film being composed of
ZrO2 and TiO2 , a highly protective thin surface film in corrosive solutions. In the NaOH solution, at
the beginning of polarization, the current values were scattered, as represented by two or three active
current peaks. Metastable pitting corrosion on surface defects, such as micropores, formed during
consolidation or the competition between the formation of active CuO and passive ZrO2 and TiO2
surface layers may have caused this scattering. For Cu60 Zr30 Ti10 BMG tested in the NaCl solution,
extremely localized pit growth occurred rapidly, leading to the breakdown of the passive film through
the galvanic corrosion mechanism.
Acknowledgments: This work was supported by the Ministry of Science and Technology of Taiwan, under grant
No. MOST 105-2221-E-019-012.
Author Contributions: Yeh-Ming Cheng and Jyun-Yu Chen carried out the sample preparation and data analysis.
Pee-Yew Lee and Chia-Jung Hu designed the experimental procedure and prepared the manuscript.
Conflicts of Interest: The authors declare no conflict of interest.
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