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Acta Materialia 79 (2014) 30–36
www.elsevier.com/locate/actamat
The ultrastable kinetic behavior of an Au-based nanoglass
J.Q. Wang a, N. Chen b,c, P. Liu c, Z. Wang a, D.V. Louzguine-Luzgin c, M.W. Chen c,
J.H. Perepezko a,⇑
a
Department of Materials Science and Engineering, University of Wisconsin-Madison, 1509 University Ave., Madison, WI 53706, USA
b
School of Materials Science and Engineering, Tsinghua University, Beijing 100084, People’s Republic of China
c
WPI-Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan
Received 22 April 2014; received in revised form 5 July 2014; accepted 6 July 2014
Abstract
Ultrastable glasses and nanoglasses are two emerging materials with novel properties that have been investigated separately. In order
to explore the combined effect of ultrastable character and a nanoglass with a nanoglobular microstructure on the kinetic behavior, the
glass transition and crystallization behaviors of an ultrastable nanoglass and a melt-spun ribbon of Au-based metallic glass were examined by differential scanning calorimetry at heating rates (/) of up to 40,000 K s1. The nanoglass shows ultrastable kinetic characters at
low / (e.g. 300 K s1) and similar kinetic behaviors at high / (e.g. 30,000 K s–1) compared to the melt-spun ribbon. The nanoglobular
interfaces remain amorphous and appear to act as a kinetic constraint to induce a higher crystallization temperature compared to the
melt-spun ribbon. The interface constraint effect disappears at 30,000 K s1. These results indicate that the nanoglobular microstructure
can act to increase metallic glass stability and provide another mechanism for the synthesis of ultrastable glass.
Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Metallic glass; Glass transition; Crystallization; Ultrastable nanoglass
1. Introduction
The competition between the formation of crystalline
and glassy phases determines the glass-forming ability during cooling of a metallic melt [1–3]. If the nucleation and
growth of crystals is limited, the glass-forming ability is
enhanced and can reach levels that allow development of
bulk metallic glass (BMG). Following synthesis, strategies
to change the properties of the glass by controlling the
glass transition phenomenon and the crystallization are
of vital importance in promoting the application of these
materials. Modification of chemical composition has been
established to be one effective way to achieve this [4], but
is limited to the glass-forming composition range. At the
⇑ Corresponding author.
E-mail address: [email protected] (J.H. Perepezko).
same time, approaches to stabilize the glass and frustrate
the nucleation and growth of crystalline phases without
changing the chemical composition are of particular
interest.
Because the free surface of glasses shows distinct properties compared to the bulk sample, such as faster diffusion
rate and stronger liquid character [5–7], vapor-deposited
glasses can show ultrastable characters with a high Tg
(glass transition temperature upon heating) if they are
deposited at appropriate temperatures, i.e. in the range of
0.7–0.95 Tg [8–11]. On the other hand, the interfaces
between the glass and substrate or between the glass and
constraining materials (i.e. porous confinements) can also
induce high stability [7,12,13]. Some polymer glasses with
nanoglobular microstructures were found to be ultrastable
[14,15]. In fact, these results suggest that the properties of
polymer glasses can be modulated distinctly by introducing
free surfaces and interfaces. Metallic glasses (MGs)
http://dx.doi.org/10.1016/j.actamat.2014.07.015
1359-6454/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
J.Q. Wang et al. / Acta Materialia 79 (2014) 30–36
represent an interesting model system to study the nature
and thermal kinetics of glasses, because of their relatively
simple atomic packing structure [16–18] and universal
kinetic behavior during the glass transition [19,20]. In addition, recently developed Fe-based nanoglass alloys have
demonstrated enhanced ferromagnetic behavior associated
with a nanoglobular microstructure [21,22], but the influence of the nanoglobular amorphous microstructure on
the thermal behavior has not been explored. By controlling
the contribution of free surfaces and interfaces [7,12–14], it
is possible to obtain nanoglobular MGs with an ultrastable
kinetic behavior similar to that observed in organic glasses.
In order to explore the combined effect of ultrastable
character and a nanoglobular microstructure on kinetic
behavior, an Au52Ag5Pd2Cu25Si10Al6 nanoglass was prepared using a slow rate (4.6 nm min1) magnetron sputtering deposition method. The thermal properties of the
nanoglass and its melt-spun ribbon counterpart were studied using high-rate differential scanning calorimetry (DSC)
which is capable of heating rates of up to 40,000 K s1. The
nanoglass shows ultrastable thermal properties, such as
20 K higher Tg and 32 K higher Tp1 (the peak temperature
of the first exothermic crystallization peak in DSC trace) at
a low heating rate of 300 K s1, compared to the ribbon
sample. However, when the heating rate increases, the differences in the thermal characteristics become smaller and
finally merge when the heating rate exceeds 20,000 K s1.
To verify the origin of the difference in the thermal kinetics,
the crystallized microstructures were examined using highresolution transmission electron microscopy (HRTEM).
2. Experimental details
The Au-based nanoglass was prepared using magnetron
sputtering with a powder target. The powder particles in
the target were high-purity elements (>99.9 wt.%) with
diameter of 100 lm. The deposition process was carried
at 300–320 K. A 10 min downtime interval was carried
out after every 10 min of continuous sputtering to minimize
heating. After six cycles (6 10 min) of sputtering the
experiment was interrupted for 30 min. After 5 h of continuous cycles of sputtering, the deposition was stopped for at
least 2 h to make sure the substrate cooled completely. The
deposition rate is 4.6 nm min1 [23,24]. To prepare the
ribbon counterpart sample, elemental components of high
purity (99.9–99.99 wt.%) were melted and flipped over four
times using an arc-melter. The ingot was then remelted in a
quartz tube in an induction-melting furnace and subsequently injected onto a spinning copper wheel (surface
velocity 30 m s1) to obtain ribbon samples. All the
above procedures were carried out under a protective
atmosphere of high-purity Ar gas.
The high-rate DSC measurements were carried out on a
Mettler Toledo Flash DSC 1 machine [25,26]. The allowable heating rates range from 10 to 40,000 K s1 and cooling rates range from 10 to 10000 K s1. The samples used
in the calorimetry measurements were cut from as-spun
31
or as-deposited samples into very small pieces under an
optical microscope and had mass ranging from 0.05 to
0.2 lg. Conventional thermal analysis was carried out on
a Perkin-Elmer DSC 8 machine at a heating rate of
0.67 K s1 with a sample mass of 5 mg.
The atomic packing structures were examined with a Cscorrected JEM-2100F transmission electron microscope.
The crystallized samples after the Flash DSC measurements were carefully transferred from the Flash DSC chip
to a TEM copper mesh with holey carbon-films for TEM
observation. Another copper mesh was pasted onto the
copper mesh with the sample encased between them to
make sure the sample was secured. When the sample is
cut from the as-deposited or as-spun sample, the wedgeshape edges can provide sufficiently thin zones for
HRTEM experiments.
3. Results
The morphology of the as-deposited MG was characterized by scanning electron microscopy (SEM), as shown in
Fig. 1. The nanoglobular microstructure of the as-deposited sample was observed in both plan view as shown in
Fig. 1a and cross-section view as shown in Fig. 1b. The statistical analysis of the globule sizes give an average value
30 ± 3 nm [23,24]. The amorphous structure for the asdeposited nanoglass and melt-spun ribbon were verified
by HRTEM and selected-area electron diffraction, as
shown in Fig. 2.
The thermal properties, including the glass transition
and crystallization of the nanoglass and ribbon sample,
were characterized by Flash DSC at heating rates ranging
from 10 to 40,000 K s1. Representative DSC traces for
the ribbon sample and nanoglass at various heating rates
are shown in Fig. 3a and b, respectively. The obvious glass
transition and crystallization phenomenon further confirm
the glassy nature of the materials. Because the samples for
Flash DSC measurements are so small (0.05–0.2 lg), it is
impossible to measure the sample weight precisely for each
measurement, so that the heat signals are not normalized
by weight. For each heating rate, 2–4 samples were measured to confirm the reproducibility (the error in Tg and
Tp1 is within 2–6 K).
The thermal characteristic temperatures including Tg
and Tp1 determined in Fig. 3 are plotted vs. heating rate
(logarithmic), as shown in Fig. 4a. The data measured by
conventional DSC at a heating rate of 0.67 K s1 (=
40 K min1) is also added for comparison. Both Tg and
Tp1 for the two samples increase with increasing heating
rate, which is typical thermal behavior for metastable
glassy materials. However, it is intriguing to find that the
nanoglass has much higher characteristic temperatures
than the ribbon sample at low heating rates. For example,
at 100 K s1 or lower, the Tg is 20 K higher and Tp1 is
32 K higher for the nanoglass compared to the melt-spun
ribbon. This denotes that the deposited nanoglass is kinetically more stable than the ribbon sample. The difference
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J.Q. Wang et al. / Acta Materialia 79 (2014) 30–36
Fig. 1. (a) Plan view SEM image of the nanoglass surface. (b) SEM image of the nanoglass cross-section. The average diameter of the nanoglobules is
30 nm.
a
ribbon
Tp1
400 450 500 550 600
T (K)
Heat flow, exo. (mW)
Heat flow, exo. (mW)
Fig. 2. (a) HRTEM image of the atomic packing structure for the as-deposited nanoglass. The inset is the selected-area electron diffraction pattern. (b)
HRTEM image of the atomic structure for melt-spun ribbon. The inset is the selected-area electron diffraction pattern. These data confirm the amorphous
atomic structure of both the nanoglass and ribbon sample.
b
nanoglass
Tp1
400 450 500 550 600
T (K)
Fig. 3. (a) Flash DSC traces of the melt-spun ribbon sample at various
heating rates, 30, 100, 300, 1000, 3000, 6000, 10,000, 20,000 K s–1 from
bottom to top, respectively. (b) Flash DSC traces of the nanoglass at
various heating rates, 10, 100, 300, 1000, 3000, 10,000, 20,000, 30,000,
40,000 K s1 from bottom to top, respectively. The peak temperatures of
the first crystallization event (Tp1) are marked by arrows. (Note: the heat
flow values are not normalized by sample weight.)
between the characteristic temperatures for the two samples becomes smaller when the heating rate increases and
finally coincides at 30,000 K s1, as shown in Fig. 4a.
To examine further the difference in thermal behavior
observed between the nanoglass and melt-spun ribbon,
the two types of MGs were pre-heated up to various temperatures above Tg (up to 25 K higher than Tg) at
300 K s1 and then cooled down immediately to room temperature at -6000 K s1. The samples were then heated up
again at 300 K s1 to measure the glass transition and
crystallization temperatures to study the effect of preannealing. The measurement sequences are illustrated schematically in Fig. S1. The DSC traces in the second heating
runs are shown in Fig. 4b for the ribbon sample and in
Fig. 4c for the nanoglass. The characteristic overshoot
behavior [8] near the glass transition temperature for the
samples disappears gradually with pre-annealing. The Tg
signals measured in the second DSC heating runs after
pre-annealing are shown in Fig. 4d. When the pre-annealing temperature increases, the Tg of the nanoglass
decreased gradually and approached the Tg of the ribbon
sample, while the Tg of the ribbon sample shows little
change upon pre-annealing. For the nanoglass, when the
pre-annealing temperature is 25 K higher (478 K) than
Tg, the Tg of the relaxed sample becomes constant and is
5 K higher than that of the ribbon sample. These results
show that the ultrastable character of the nanoglass can be
completely removed by heating to higher temperatures at
J.Q. Wang et al. / Acta Materialia 79 (2014) 30–36
a
Tp1-nanoglass
Tp1-ribbon
550
Tg-nanoglass
T (K)
Tg-ribbon
500
450
400
0
1
2
3
4
5
Heat flow endo. (a.u.)
Tg
Tg
ribbon
as-spin
433 K
438 K
448 K
Heating rate: 300 K/s
400 420 440 460 480
T (K)
10 10 10 10 10 10
Heating rate (K/s)
c
b
Heat flow, exo. (mW)
600
460
nanoglass
d
as-deposit
450
Tg (K)
453 K
463 K
468 K
473 K
Heating rate: 300 K/s
400
as-deposit
450
500
T (K)
550
453 K
463 K
nanoglass
468 K 473 K
440
478 K
430 as-spin433 K 438K 443 K448 K 451 K
ribbon
420
1 2 3 4 5 6
sample no.
Fig. 4. (a) Tg (solid symbols) and Tp1 (open symbols) for both nanoglass
(squares) and ribbon (triangles) vs. heating rate. DSC traces of (b) ribbon
and (c) nanoglass after pre-annealing to various temperatures. The preannealing temperatures are marked for each curve. (d) Change in Tg for
ultrastable nanoglass (squares) and ribbon (triangles) after pre-annealing
at different temperatures which are marked at each point.
0
a
300
Δ E (kJ/mol)
ribbon
2
ln(φ/Tp1 )
-3
-6
nanoglass
-9
b
200
nanoglass
100
ribbon
-12
1.6
1.8
2.0
2.2
-1
1000/Tp1 (K )
2.4
0
1.6
1.8
2.0
2.2
-1
1000/T (K )
Fig. 5. (a) Kissinger plot of the crystallization peak in DSC traces for both
nanoglass and ribbon. (b) Crystallization activation energy determined
from the slope of the Kissinger plots.
an ultra-fast heating rate or partially removed by pre-heating at a low heating rate to the supercooled liquid state.
The crystallization activation energy (DE) derived from
the slope of the Kissinger plot [27] of ln(//T2p1) vs. 1000/
Tp1 (see Fig. 5) is approximately equal to that for the
growth of polymorphic crystallization [28]. As shown in
Fig. 5b, the DE for the nanoglass is larger than that for
the ribbon sample, especially at low temperatures (at low
heating rates), which is additional evidence for the ultrastable nature of the nanoglass. The crystal growth rates (U)
were estimated experimentally by measuring the average
crystal size (radius, r) and the crystallization time (t) using
HRTEM and DSC, with U ¼ r=t (see Fig. 6). The experimentally determined growth rates are shown as symbols
in Fig. 7. The crystal growth rate can also be estimated
33
from the Kissinger plot according to Johnson–Mehl–
Avrami (JMA) [29] kinetics and is given by the solid curves
shown in Fig. 7 (details can be found in the online Supplementary Material and Fig. S2). The theoretical evaluation
is consistent with the experimental estimation. This shows
that the ultrastable nanoglass has a lower crystal growth
rate at low T (or at low heating rates) than that for the ribbon sample.
4. Discussion
The temperature of the substrate in fabricating the
nanoglass is between 300 and 320 K (considering a temperature fluctuation of 20 K due to the exothermic plasma
deposition process), which represents 0.75Tg–0.8Tg, given
that the Tg of Au52Ag5Pd2Cu25Si10Al6 melt-spun ribbon
is 400 K (at 0.67 K s1) [24]. With the slow deposition
rate (4.6 nm min1) and the high substrate temperature,
the atoms are able to relax to stable states with low energy
and high packing efficiency [8–11]. Because sufficient free
volume is required for the glass transition [30], the asdeposited Au-based metallic glass requires a higher temperature than the melt-spun ribbons to expand to generate
enough free volume upon heating to exhibit a glass transition. These results are similar to experimental observations
on vapor-deposited organic glass formers that demonstrate
that an ultrastable state can be achieved if the substrate
temperature during deposition is between 0.7 and 0.95Tg
[31].
The formation of a nanoglobular microstructure has
been widely observed in depositing both liquids and metals
on substrates [32–36]. For liquid deposition (i.e. water), the
homogeneous formation of hemispherical droplets on a
substrate is well known as a so-called “breath figure”
[32–34]. For pure metal deposition (i.e. Ag and Cu), the
surface atomic diffusion allows atoms to nucleate into
islands. The globular islands can be retained even after they
grow and contact each other if the diffusivity is low enough
to avoid coalescence [35,36]. For a glass-forming system,
the droplets formed upon deposition on a sufficiently cold
substrate would become a glass because of the fast cooling
rate. If the diffusivity is low enough to prevent the coalescence at the substrate temperature, the nanoglobular
microstructure can be retained, which is similar to the
formation of crystalline nanograins [37].
Based upon the contrast variation in TEM images
(Fig. 8) and reports in the literature [23,24,37], the nanoglobular interfaces exhibit a lower density than the interior.
While the lower density is usually associated with enhanced
diffusive transport, the nanocrystallization microstructure
reveals that the nanocrystals initiate within the interior.
At the same time, there is evidence (see Fig. 8a) that the
nanoglobular interfaces are retained as amorphous regions
even after the interior has transformed to several nanocrystals, which is verified by the fast Fourier transformation
(FFT) images in the insets of Fig. 8a. This indicates that
the nanoglobular interfaces do not act as preferred sites
34
J.Q. Wang et al. / Acta Materialia 79 (2014) 30–36
-3
Growth rate (m/s)
10
-5
10
-7
10
ribbon
nanoglass
-9
10
1.6
1.8
2.0
2.2
-1
1000/T (K )
2.4
Fig. 7. The measured and simulated growth rate of crystals for both
nanoglass and ribbon vs. reciprocal temperature.
Fig. 6. (a) TEM image of the ribbon crystallized isothermally at 443 K.
The crystal diameter is 36 nm. (b) The corresponding DSC trace shows
that the crystallization lasts for 0.9 s. The crystal growth rate is
20 nm s1. (c) TEM image of the ribbon crystallized with continuous
heating at 100 K s1. The average crystal diameter is 40 nm. (d) The
corresponding DSC trace shows that the crystallization lasts for 0.06 s
with a peak temperature of 456 K. The crystal growth rate is
333 nm s1. (e) TEM image of the nanoglass crystallized with continuous heating at 100 K s1. The average crystal diameter is 8 nm. (f) The
corresponding DSC trace shows that the crystallization lasts for 0.18 s
with a peak temperature of 488 K. The crystal growth rate is
22 nm s1. (g) TEM image of the nanoglass crystallized with continuous
heating at 30,000 K s1. The average crystal diameter is 50 nm. (h) The
corresponding DSC trace shows that the crystallization lasts for 1.8 ms
with a peak temperature of 564 K. The crystal growth rate is
14 lm s1.
for nanocrystal nucleation and appear to act as kinetic barriers to nanocrystal growth. The origin of this intriguing
behavior requires further study to elucidate the mechanism.
It has been also reported that the interface between the
glass and substrate or the constraint by porous materials
can enhance glass stability [12,13], and that some nanoglobular polymer glasses can show ultrastable character
[14]. This suggests that the nanostructures and their interfaces in the deposited glass could contribute to the ultrastable nature. However, when the sample is heated to above
the crystallization temperature at a high heating rate
(30,000 K s–1), the kinetic constraint is removed since both
the interior and the interface are crystallized and a crystal
can grow across the nanoglobular interfaces, which is also
verified by the inset FFT image in Fig. 8b.
Pre-heating at a low heating rate (300 K s1) to slightly
above Tg can partially remove the ultrastable character by
releasing the relaxation effect (because of the slow deposition rate in fabricating the sample) within the nanoglass
but cannot remove the interface constraint, while heating
at ultra-fast heating rates can completely remove the
ultrastable character by releasing both the relaxation effect
and the constraint of the nanoglobular interfaces.
An important signature of ultrastable glasses is the characteristic lower enthalpy and associated reduced volume
compared to ordinary glasses [8–11]. One consequence of
these characteristic features that is directly related to the
observed behavior is reduced diffusivity. The lowered diffusivity delays the onset of kinetic reactions upon heating.
With the same crystallization reaction, the lower diffusivity
of an ultrastable glass will result in an increase in the transient period before the onset of steady-state nucleation [38].
As a result, the Tp1 on heating for an ultrastable glass will
be higher than that for an ordinary glass. However, since
the diffusivity is temperature dependent, the transient period will shorten at higher Tp1 and eventually merge with
the Tp1 for an ordinary glass (see Fig. 4a).
For ultrastable glasses, the stable kinetic characters (e.g.
higher glass transition temperature) are usually related to
their stable thermodynamic properties (e.g. low energy
states or low enthalpy) [8–11,39]. However, recent research
has shown that ultrastable kinetic behavior was also found
J.Q. Wang et al. / Acta Materialia 79 (2014) 30–36
35
Fig. 8. (a) HRTEM image near the nanoglobular interface confirming that the interface remains amorphous after heating to above the DSC
crystallization event at 300 K s1. The upper-right inset is the FFT image of the nanoglobular interior verifying the crystalline atomic structure, and the
lower-right inset is the FFT image of the interface verifying the amorphous atomic structure. (b) HRTEM image near the interface confirming that the
crystal can grow across the interface after heating to above the crystallization event at 30,000 K s1. The inset is the FFT image of the interface verifying
the crystalline atomic structure.
Enthalpy (J/g)
5. Conclusion
5 J/g
Ribbon:
preannealed at 438 K
Ribbon:
as-spun
Nanoglass:
preannealed at 468 K
Nanoglass: as-deposited
370
400
430
T (K)
460
490
Fig. 9. The enthalpy of the as-deposited nanoglass, the nanoglass preannealed at 468 K, the as-spun ribbon, and the ribbon pre-annealed at
438 K. The enthalpy of the nanoglass changes by 1.0 J g1 after preannealing at 468 K. The enthalpy difference between the as-deposited
nanoglass and as-spun ribbon is 1.5 J g1. The enthalpy of the ribbon
changes <0.2 J g1 after pre-annealing at 438 K. The solid grey line shows
the extrapolation of the enthalpy for the supercooled liquid.
in a metallic glass [40] and a polymer nanoglass [14] with a
higher enthalpy than ordinary glass. The mass of the Flash
DSC sample (<10 lg) cannot be measured using a balance.
It was estimated based on the crystallization enthalpy. The
enthalpy of the as-deposited nanoglass is about 1.5 J g1
lower than that of the as-spun ribbon (see Fig. 9). After
pre-annealing at 468 K, the enthalpy of the nanoglass sample increases by 1.0 J g1 due to the loss of its ultrastable
character. The fictive temperature (Tf, the cross-temperature of the glass enthalpy and extrapolation of the supercooled liquid enthalpy) for the as-deposited nanoglass is
384 K (41 K lower than that of the as-spun ribbon,
Tf = 425 K) and increases by 23–407 K after pre-annealing.
Thus, the Au-based nanoglass is also thermodynamically
ultrastable.
In summary, deposited Au-based nanoglass shows both
kinetic and thermodynamic ultrastable characters at low
heating rates compared to its melt-spun ribbon counterpart. The nanoglass shows a 20 K higher glass transition
temperature (Tg), a 32 K higher crystallization peak temperature (Tp1), larger crystallization activation energy,
slower crystal growth rate, and lower enthalpy compared
to the ribbon sample. The ultrastable behavior is attributed
to both the relaxation during deposition and the nanoglobular microstructure. The ultrastable character can be partially removed by pre-heating up to supercooled liquid
region at low / and can be completely removed by heating
to higher temperature at extra-high /. These results indicate that a glass can be made more stable both by relaxation during fabrication and by synthesis of a
nanoglobular microstructure. This provides new opportunities to modify the stability and properties of metallic
glasses.
Acknowledgements
We thank Professor M.D. Ediger, University of Wisconsin–Madison for valuable and insightful discussions. The
financial support from the NSF (DMR-1005334), ONR
(N00014-12-1-0569) and World Premier International
Research Center Initiative (WPI)-MEXT Japan are gratefully acknowledged.
Appendix A. Supplementary data
Supplementary data associated with this article can be
found, in the online version, at http://dx.doi.org/10.1016/
j.actamat.2014.07.015.
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J.Q. Wang et al. / Acta Materialia 79 (2014) 30–36
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