of 10
Nano
-
metallic glasses: size reduction brings ductility, surface state drives its extent
D.Z. Chen
*,1
, D. Jang
1
, K.M. Guan
2
, Q. An
3
, W.A. Goddard,
III
3
, and J.R. Greer
1,4
,
1
Division of Engineering and Applied Sciences, California Institute of Technology, Pasadena
CA, 91106
2
Department of Chemistry and Chemical Engineering, California Institute of Technology,
Pasadena CA, 91106
3
Materials and Proces
s Simulation Center, California Institute of Technology, Pasadena CA,
91125
4
The Kavli Nanoscience Institute, California Institute of Technology, Pasadena CA, 91106
1.
Supporting
Figures:
Figure S1: Average Young's moduli for FIB and EP samples across ten independent specimens
(five of each type)
. EP samples show a slightly lower modulus, but it is within one standard
deviation.
Figure S2
: Graph
depicting the d
2
versus d
3
scaling argument for surface energy, E
Surface
and total
stored elastic energy, E
Elastic
. At diameters lower than d* (opaque green region), the energy
required for a crack/fracture to initiate at the surface (E
Surface
) dominates,
and shear banding is
suppressed. Altering the surface state increases the gap between E
Surface
and E
Elastic
, further
suppressing catastrophic failure.
Figure S3
Diagram depicting the generation of high potential energy atoms from ion irradiation
for a
typical binary MG system. Top: as
-
cast configuration of NiP MG atoms, with the free
surface to the right of the depicted atoms (not shown). Ga ions impinge upon the free surface
from the right. Bottom: atoms rearrange after collision, resulting in high pot
ential energy
atoms
shown in light green for Ni and light blue for P. The free volume in the vicinity of these
atoms is notably higher, creating an easy path for them to take in response to elastic energy
input.
2.
Supporting
Methods
Synthesis
Electroplated (EP) samples were fabricated following a template pattern transfer
procedure described in.
1
In this methodology, a thin layer of Au (~100 nm) is first evaporated
onto a silicon wafer coated with a 3
0nm
-
thick Ti layer without breaking the vacuum. The
~700nm
-
thick PMMA is then spin
-
coated onto this seed layer, patterned using e
-
beam
lithography, and developed to reveal vertical through
-
holes with the desired diameters. The
electroplating is performed b
y providing a constant current between the anode (Ni foil) and the
cathode (prepared template and Au dummy chip), and tension samples are made by overplating
the metallic glass above the PMMA surface. Table 2 provides details on the specific Ni
-
P
electropl
ating conditions. More details about this fabrication method for other metallic nano
structures can be found in
1
.
In addition to such templated electroplating, a blanket ~2μm
-
thick Ni
-
P film was
electroplated di
rectly onto the seed layer as a separate Si chip. This film had virtually identical
composition to the nano
-
pillars (Table 2 and Figure 2f). Tensile samples with ~100
-
nm and 500
-
nm diameters and ~650
-
nm and 2
-
μm gauge lengths were then fabricated from an e
lectroplated
NiP film using an FEI Nova 200 Nanolab focused ion beam with a final etching condition of
30kV/10pA (Figure 3). Great care was taken when carving the heads of the tension samples to
minimize bending during nano
-
mechanical testing.
Tensile Te
sting
A custom
-
made
in situ
SEM, SEMentor (SEM + nanoindenter), was used to conduct the
nanomechanical tension tests.
2
The nanoindentor arm is fitted with a diamond Berkovich tip that
was carefully milled with FIB into the shape and dimensions of a tensile grip, as seen partially in
Figure 3. All of the experiments were carried out
at a constant nominal displacement rate (0.5
-
8
nms
-
1
) using a feedback algorithm, which results in a global strain rate of 1.0e
-
3
s
-
1
. The raw load
-
displacement data were recorded after isolating the specimen
-
only response from the load frame,
support sprin
g, and substrate compliances. Engineering stress and strain values were then
calculated from load
-
displacement data using the sample diameters measured from SEM images.
Molecular Dynamics
We started with a Ni
3
Al crystal with a total of 4000 atoms and melt
ed the system at 3000
K. Then 5% of the Al atoms were replaced by Ni atoms to form a Ni
80
Al
20
liquid.
Periodic
boundary conditions were applied in all three directions of the simulation box to eliminate
surface effects. Equations of motion were solved usin
g the velocity Verlet algorithm with a time
step of 1 femtosecond. The dynamics were carried out with an NPT ensemble (constant particles,
pressure, and temperature) using a Nose
-
Hoover thermostat (time constant of 0.1 picosecond)
and barostat (time consta
nt of 1 picosecond).
In addition to the steps mentioned in the main text of the paper, periodic boundary
conditions were applied only along the cylinder axis during quenching. Additionally, the length
of the cylinder was fixed during this step as well. Th
e irradiation step was performed using an
NVE ensemble (constant particles, volume, and energy) for 200 picoseconds until the potential
energy of the system stabilized. After this the irradiated pillars were immediately quenched to
room temperature and rel
axed using an NVT ensemble (constant particles, volume, and
temperature) for 500 picoseconds. To avoid an unphysical attraction in this high
-
energy collision
process, we applied a truncated and shifted Leonard
-
Jones pairwise interaction to the original
EAM
potential in the repulsive region. The form and parameters of the LJ potential are identical
to that of Xiao, et al.
3
Uniaxial tension was conducted by rescaling the simulation box along the
loading direction. The atomic stresses were obtained from the atomic virial
4
to extract the total
stresses of the nanopillars.
The simulated irr
adiation fluence value of 0.0625
/nm
2
was obtained from a back
-
of
-
the
-
envelope calculation using experimental FI
B conditions on the FEI Nova 200 NanoLab
DualBeam
TM
SEM/FIB. The dwell time for the beam was 1
μs, and
the spot size was 2 nm
2
.
From these two parameters, we can
use eq. (1) to
estimate the number of scans per milling
session, given the user
-
specified FIB
pattern area and a
correction to the scan area for the 50%
overlap in beam spot in both the x and y directions.
From this estimate of the number of scans
per session, we can
use eq. (2) to
obtain the ion fluence
by applying the number of ions per
second needed to achieve an ion current of 10 pA, whi
ch results in a value of ~0.0625
/nm
2
.
(1)
*
*
:
T
he factor
2 corresponds to a correction to the scan area for the 50% overlap in the beam spot for the x and y
directions.
A
Spot
is the spot size, A
FIB
is the total pattern area, t
Final
is the final exposure time, and t
Dwell
is the total
dwell time.
(2)
**
**
:
The factor 10 corresponds to a rough correction for the glancing angle of incidence for the incoming ion beam.
I
Ion
is the ion current, t
Sample
is the FIB exposure time on the sample surface, e is the electron charge, and A
FIB
-
afffected
is the cross
-
sectional area of the FIB
-
affected zone on the sample surface.
3.
Supporting
Discussion:
Short range order (SRO)
A Ni
-
P system was chosen because
it lends itself well to electroplating. NiP metallic
glasses may have different short range order (SRO) compared with the more common binary
glass, CuZr, which has more metallic
-
like bonding.
5
Although the bonding between Ni and P
may not be purely metallic, it should not adversely affect the role of Ni
-
P metallic glass as a
model system nor does it change the findings or the
conclusions of this work.
Neutron diffraction
experiments have also suggested that the pair distribution functions (PDFs) of metal
-
metalloid
systems are qualitatively similar to metal
-
metal systems such as CuZr, especially after the second
nearest neighbor
s.
6
It is not known
whether the brittle
-
to
-
ductile transition occurs in only
metallic glass systems or in other amorphous systems as well. A similar phenomenon has been
seen in amorphous silica nanowires, which at 50
-
100nm diameters undergo brittle failure
7
and at
down to 20nm exhibit great deformability
8
. The Cu
-
Zr which was shown to undergo the size
-
induced brittle
-
to
-
ductile transition
(ref. 10 in the manuscript) has more of a metallic
-
type
bonding, but cannot be considered to be purely metallic. Theoretically, the glass formability of a
purely metallic bonding system is basically zero
for example a slab of pure Cu or pure Ni
cannot be
made to be amorphous. Typically, a certain amount of covalent bonding or orientation
preference is needed for the formation of an amorphous microstructure. Additionally, the
microstructure of our NiP system is completely amorphous, as validated by TEM, and
, as Argon
et al.’s (ref. 25 in the manuscript) experiments have shown, even unrelated analogous systems
such as bubble rafts can lead to rather applicable insights to the underlying deformation in
amorphous systems, including metallic glasses.
Electroplating
The rate of metal ion deposition during electroplating was fairly consistent for both the
films and the 100
-
nm template pillars, suggesting that the plating time is independent of template
feature size. Even at 100nm
-
diameters, the electropl
ating remained reasonably unhindered by the
template, which implies that the only relevant driving force for plating rate is the current density.
From the SEM images, it can also be seen in Figure 3
-
b & c that the electroplated pillars had
relatively smoot
h surfaces, with few hydrogen bubbles, which are a common issue with
electroplated samples.
9
Tensile specimens also have characteristic caps, which are a result of
plating over the e
-
beam developed features. These caps are quite isotropic, and somewhat
resemble mushroom tops. Isotropy in the overplated caps is an i
ndicator for the presence of an
amorphous rather than a crystalline phase, which might show some sort of anisotropy or
orientation preference, as we have seen with our previous samples.
1
Carbon and Oxygen contamination
EDX only detects the surface atoms, and due to th
e small size of the EP specimen, the
spot size of the e
-
beam actually includes some of the atoms on the template surface as well as
those on the specimens. We have re
-
normalized the chemical compositions based only on the
relative Ni and P content to refle
ct this. The resulting difference in relative P content is less than
1 wt% between the two sample types, which makes them be nominally identical in terms of
chemical composition, as planned. As mentioned in the manuscript, the O and C present on the
surfac
es of the samples should not form continuous layers, and as such do not bear any load. It
may affect the straining slightly, but it is highly unlikely to induce plastic flow in an otherwise
brittle glass.
Strain from in
-
situ video
Raw data obtained from
the displacement signal of the nanoindenter produces results
which have high variability in the loading moduli, an unavoidable characteristic detriment of
nano
-
mechanical testing.
Many factors can affect the apparent modulus: slight misalignments
between t
he sample and the grip, surface imperfections on the sample or the grip (bumps), the
machine compliance, the imperfect adhesion, etc. to name a few. We have carefully re
-
examined
the raw data and used the frames from the in
-
situ SEM videos rather than the
displacement signal
from the nanoindenter to measure and recalibrate the strains obtained in the mechanical data.
This technique is more reliable for strain determination and has been successfully utilized in
many publications from our group.
10
-
12
Such a discrepancy between the strain based on video
frames and that measured in the nanoindenter is typical and likely stems from the
effects of
machine and grip compliances in both sample types and additional substrate effects in EP
samples.
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