Extreme Mechanics Letters 13 (2017) 126–134

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Extreme Mechanics Lettersjournal homepage: www.elsevier.com/locate/eml

Liquid metal particle popping: Macroscale to nanoscale
Trevor R. Lear a , Seok-Hee Hyun b , John William Boley a , Edward L. White a ,David H. Thompson b , Rebecca K. Kramer a,∗aSchool of Mechanical Engineering, Purdue University, 585 Purdue Mall, West Lafayette, IN 47907, USAb

Department of Chemistry, Purdue University, 560 Oval Drive, West Lafayette, IN 47907, USA

articleinfo

Article history:
Received 18 November 2016
Available online 28 February 2017
Keywords:
Gallium–indium alloy
Liquid metal
Nanoparticles
Atomic force microscopy
Particle rupture
Nanoindentation

abstract
Liquid metal nanoparticles can be used to produce stretchable electronic devices. Understanding
the mechanical properties of liquid metal nanoparticles is crucial to optimizing their use in various
applications, especially printing of flexible, stretchable electronics. Smaller nanoparticles are desiredfor high-resolution printing and compatibility with existing scalable manufacturing methods; however,they contain less liquid metal and are more difficult to rupture than larger particles, making themless desirable for post-processing functionality. This study investigates the mechanics of liquid metal
particle rupture as a function of particle size. We employ compression of particle films to characterize
the composition of the particle core and derive a minimum particle size required to achieve sintering andsubsequent conductance. We further derive the force required to rupture a single particle and validate
the results by rupturing individual nanoparticles using atomic force microscopy. Finally, we relate theliquid metal nanoparticles to isotropically-elastic thin-shell microspheres to approximate the particle
shell stiffness. An increased understanding of the behavior of liquid metal nanoparticles during rupture
reveals limitations of current manufacturing processes and paves the way for the next generation ofscalable mass-producible soft electronics using additive manufacturing technologies.
© 2017 Elsevier Ltd. All rights reserved.

1. Introduction
Liquid metals offer new opportunities for flexible, stretchable,
and shape changing electrical components [1–10]. Effective techniques to process liquid metal have been demonstrated and include injection into microchannels [1,3], imprinting [11], masked
deposition [12], and extrusion [13–15]. Although it is possible
to manipulate liquid metals at submillimeter length scales, thesetechniques are greatly inhibited by the spontaneous formation of athin metal oxide layer on the liquid metal surface in the presence ofoxygen. This metal oxide is the mechanism behind the unique capability of liquid metals to form free-standing structures [14,16],
but also produces a high surface tension that makes them incompatible with scalable liquid processing techniques [16,17], such as
inkjet printing. Inkjet printing is desirable due to its capacity to create high resolution patterned devices while remaining a high-yieldprocess. The ability to inkjet print any liquid for a particular application is indicated by the Ohnesorge number, which relates the∗ Corresponding author.
E-mail address: rebeccakramer@purdue.edu (R.K. Kramer).
http://dx.doi.org/10.1016/j.eml.2017.02.009
2352-4316/© 2017 Elsevier Ltd. All rights reserved.

viscous forces to inertial and surface tension forces, and liquid
metal has been shown to be outside the range of printability [18].
We have recently demonstrated the ability to inkjet-print
liquid metals by dispersing liquid metal nanoparticles in a
printable carrier solvent, which allows us to leverage thecarrier solvent properties during processing and the liquid metal
properties post-deposition [19]. The liquid metal nanoparticles
are electrically insulated by an oxide shell that may be ruptured
via application of pressure to sinter the particles and form anelectrical path. However, optimal particle preparation conditions
for compatibility with drop-on-demand printing and functionality
post-processing are contrasting: particle size should be minimized
for compatibility with printing to avoid nozzle clogging, while
particle size should be maximized for electrical functionality toincrease the ratio between conductive liquid metal and semiconductive metal oxide after particle sintering.
Here, a detailed characterization of liquid metal particle
sintering through particle rupture is described. We characterize
the electrical response of liquid metal nanoparticle films using
compression testing and derive a minimum particle size required
to achieve sintering and subsequent sheet conductance. We further
derive the force required to rupture a single particle as a functionof particle size and introduce a composition model for the liquid

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metal particle core and corresponding metal oxide shell thickness.
These results are validated using an atomic force microscopy (AFM)
nanoindentation method, through which the individual particles
are ruptured with an AFM tip. This work suggests that at submicron scales the repeated formation of the metal oxide shellaround the liquid metal particles may greatly alter the compositionof the particle core. The inherent relationships between particle
formation, core and shell composition, and conductance after
coalescence are expected to enable scalable manufacturing of
liquid metal-based soft electronics.
2. Nanoparticle formation and characterization
In an approach similar to that of Hohman et al. [17] and revised
by Boley et al. [19], we created stable ethanolic nanoscale liquid
metal colloids using sonication, which induces mechanical separation to form liquid metal nanoparticles due to the presence ofhigh oscillating shear forces, as shown in Fig. 1(a). The liquid metal
employed in this work is a commercially available eutectic gallium–indium (EGaIn; 78% Ga and 21.4% In by mass; 15.5 °C melting point, Figure S1). Due to exposure to oxygen during the sonication process, gallium oxide continuously forms, fractures, andreforms, which results in decreasing particle size with increasing
sonication time (Fig. 1(b)). Similar to the gallium oxide that forms
on pure gallium particles [20], the semi-solid oxide layer is electrically insulating and mechanically stabilizing, which prevents spontaneous particle coalescence [21–24]. After sonication, a broad distribution of particle geometries were observed [25,26]. Previous
experiments have employed thiols (for example, 3-mercapto-Nnonylpropionamide (1ATC9)) and polyvinylpyrrolidone (PVP) toachieve geometric and size uniformity, as well as encourage propersuspension [17,19]. Here, we examine liquid metal nanoparticles
with oxide coatings and thiol coatings to characterize the mechanical and electrical properties of both types of particles.
Using identical methods to Boley et al. [19], we produced liquid
metal nanoparticles both with and without thiols with sonication
times ranging from 30 to 960 min. The particle sizes werecharacterized using scanning electron microscopy (SEM), withselected samples shown in Fig. 1(c) (SEM images for all samples
used can be found in the Supporting Information). Previous results
indicated the average particle size started to plateau at slightly
above 200 nm for particles without stabilization agents (thiols);
however, those studies did not explore sonication times beyond
120 min [19]. Samples generated by significantly longer sonication
times returned a much smaller asymptotic average particle sizeof approximately 50 nm, as Fig. 1(b) shows for thiolated andnon-thiolated particles. As seen in Fig. 1(c), samples sonicated forshorter periods of time (e.g., 30, 90 min) display a larger variancein particle size, with average diameters decreasing with increasing
sonication time. It can also be seen that the presence of thiols
in solution aids in the formation of smaller, more monodisperse
particles; the particle size distribution in non-thiolated thiolated
samples shown in Fig. 1(d) is much greater than that of thiolated
samples shown in Fig. 1(e).
3. Particle rupture experiments and results
3.1. Particle film rupture
Particle film compression tests were used to determine therelationship between mean particle size and the pressure required
to sinter the particles. Mechanical sintering of a film of liquid metal
particles has been previously studied [19], and here we extend thiswork to include both thiolated and non-thiolated particles overa greater range of sizes. Using a single column materials tester,
thin films of liquid metal nanoparticles were compressed using

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two parallel plates while simultaneously measuring the sheet
resistance, as shown in Fig. 2(a). The rupture forces of the particles
found using these experiments are represented in Fig. 2(b) which
compares the approximate rupture force per particle versus theaverage particle diameter. Fig. 2(c)–(d) show raw data for boththe non-thiolated and thiolated particles. The curves consistently
exhibit the following pattern as the particle films are compressed;(1) a non-conductive state as the flat plate contacts and begins
compression of the film, (2) a jump in voltage that signifies
successful sintering of the particle film, (3) a continued conductive
phase as the film is further compressed, and (4) a saturation stateat which compression forces ink out from underneath the plates
and breaks the conductive path. Some noise is visible in the data,
particularly during the conductive phase; we expect this is due toa combination of polydispersion of the particles and the complex
load distributions between particles, both of which may lead tosome particles in the film rupturing before others.
As seen in Fig. 2(c)–(d), although the thiolated and nonthiolated particles were sonicated for identical durations, the thiolated particles were smaller and more monodisperse, and therefore
required greater forces to achieve conductance. We also note thatdata for films comprised of smaller particles (d < 70 nm) are notavailable, as those particle films did not become conductive over
the range of forces applied. Using data from these experiments, a
relationship between particle size and the required force for sintering was developed (Fig. 2(b)). As detailed in previous work [19],the approximate rupture force per particle, f , can be calculated as:
f =

Fd2
pD2

(1)

where F is the applied force required to sinter the particle film,
D is the diameter of the flat polymethyl methacrylate punch used
to compress the particles (4.35 ± 0.3 mm), d is the average particle diameter, and p is the areal packing factor for the particles(assumed to be a constant 0.82 [27]). The resulting rupture force
per particle as a function of particle diameter shows linear dependence, which agrees with similar studies focused on rupture of therigid shell around fluid-filled microcapsules [28,29]. The smaller
variance in particle size for smaller mean particle diameters is reflected in Fig. 2(b), as the error bars decrease with decreasing particle size. Fig. 2(b) also shows that the particle rupture force forthiolated and non-thiolated particles is similar for small particles
and diverges with increasing particle size. This implies that the mechanical properties of the thiol coating and the metal-oxide coating
converge at small particle sizes (<100 nm).
3.2. Particle core composition
Let us consider a single gallium–indium liquid metal nanoparticle under sonication. We assume that the initial composition of
the nanoparticle is 78.6% Ga and 21.4% In by mass (determined by
the eutectic point of gallium–indium alloy, phase diagram shownin Figure S1), that the mass of the metal is conserved, and that gallium oxide formation on the surface of the nanoparticle contains no
trace indium. As gallium oxide forms on the surface of the nanoparticle, gallium composition in the particle core decreases and thecore becomes indium-enriched. Fig. 3(a) shows the relationshipbetween particle size, gallium oxide thickness, and the percent
of gallium in the particle core. We note that the gallium–indium
phase diagram shows that a gallium–indium alloy will no longer
be liquid at room temperature for gallium compositions less than
73%. Therefore, our model indicates a threshold at which the particle core composition will transition from liquid to solid at room
temperature. As noted previously, we were unable to experimentally sinter particles with diameters less than 70 nm. Inspecting

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Fig. 1. (a) Schematic showing the nanoparticle fabrication process. (b) Average particle diameters for thiolated and non-thiolated dispersions as a function of sonication
time. Error bars shown represent one standard deviation. (c) SEM images of non-thiolated and thiolated particles at various sonication times. Scale bar represents 2 µm.
(d)(e) Histograms showing a large polydisperse distribution of particle sizes for the 30 min sonication samples and significantly narrower size distribution for samples
sonicated for 960 min (n = 1000) for non-thiolated and thiolated particles, respectively.

our model at this minimum rupture threshold, we see that the liquid–solid core transition for 70 nm diameter particles corresponds
to an oxide thickness of approximately 3 nm. Multiple studies haveattempted to quantify the thickness of the gallium oxide shell using
X-ray scattering and/or photoelectron spectroscopy, resulting in arange from ∼0.5 to 5 nm [20,30,31]. Lin et al. and Ren et al. recently
performed studies using TEM and quantified the oxide thicknessof liquid metal nanoparticles at 3 nm for similarly sized particles[32,33], which supports the prediction of our model. Our model
further predicts that liquid metal particles with diameters less than
70 nm and an oxide thickness of 3 nm should have solid cores,
which aligns with our previous observation that particle films comprised of small particles (d < 70 nm) cannot be sintered at room
temperature.
We validated our model via compression experiments using
particles that were approximately 55 nm in diameter, which is lessthan the predicted 70 nm threshold that the model suggests willproduce solid-core particles under these fabrication conditions.
We compressed (sintered) particle films in a materials testerat ambient temperature, where the particle core is predicted
to be solid, and an elevated temperature above the melting
point of the solid gallium–indium alloy core. For the elevatedtemperature tests, the particle films were deposited onto a plate
heated to 50 °C and allowed to sit for approximately 10 min toallow the temperature of the films to stabilize. The films werethen compressed using the same parameters as the original film
compression tests. As visible in Fig. 3(b), heated particle films weresuccessfully sintered by applying between ∼300 and 350 N offorce, whereas particle films at ambient temperatures continued
to record no electrical response. This result is consistent with our
hypothesis that long sonication times result in small and solid-core
particles that cannot be ruptured in ambient conditions.

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Fig. 2. (a) Compression testing of liquid metal nanoparticles using a materials tester. Liquid metal nanoparticle dispersion (50 µL) is deposited into a PDMS reservoir fittedwith copper wires to measure the resistance during compression. The coalescence of particles is represented in the inset, where particle packing orientation and density are
not always known. (b) Particle film compression data—average rupture force per particle vs. mean particle diameter. Error bars represent one standard deviation (n = 5).
Linear approximation shown. (c)(d) Compression experiment data for non-thiolated and thiolated particles, respectively.

Fig. 3. (a) Model relating particle size, gallium oxide thickness, and percent gallium composition in the particle core. Blue region represents particles with liquid cores(which may be ruptured), pink region represents particles with solid cores (which cannot be ruptured). Equation of approximate phase transition line is D = 20.0072t
where D is the particle diameter (nm) and t is the oxide shell thickness (nm). (b) Results of film compression testing of ∼55 nm particles during heating, activation only
achieved for samples heated during compression (n = 3). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of thisarticle.)3.3. Single particle rupture
In order to validate the rupture force per particle as derived
from the bulk particle film compression data, we sought to conduct rupture experiments on individual nanoparticles using atomic
force microscopy (AFM). To isolate single particles, we first created a monolayer of nanoparticles using the Langmuir–Schaefer
method, as seen in Fig. 4(a) (process detailed in Figure S4) [34–37].

Similar experiments have shown that AFM compression of microcapsule monolayers does not differ from compression of isolated
microcapsules [38]. Because uniform particle size and geometry
are desired to produce a monolayer, thiolated particles were used
to encourage monodispersion [17]. All particles were sonicated forthe same duration and underwent differential centrifugation toeliminate extreme variation in particle size, aiding in the formationof a relatively uniform particle monolayer. Fig. 4(b)–(c) show AFM

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height and phase maps that confirm the popping of an individualliquid metal nanoparticle, which was further verified using SEM, atechnique that has previously been used to identify the rupture offilled particles [39,40].
A single particle rupture experiment is shown in Fig. 4(d)–(e)and comprises of the following sequence: (1) The AFM tip (∼5 nm
at the thickest point) is lowered until contact is made with theparticle shell. (2) Variation in force can be seen in the early regionsof contact before the particle shell fractures, which is consistent
with previous AFM indentation experiments on hard-shelled fluidfilled capsules that have shown buckling, bulging, and dimplingof the outer shell as the localized pressure increases [33,41–43].
(3) The tip completely penetrates through the shell, at which
point the particles are assumed to be ruptured. (4) Following
penetration, complete fracture of the shell occurs, allowing liquid
to flow out. (5) Finally, as the liquid flows from the fractured
particle shell, contact with ambient oxygen causes an oxide
to form, fracture, and reform under the AFM tip, as indicatedby the force oscillations shown in Fig. 4(d). We note that thissingle particle rupture process includes both the compressive andadhesive force as the AFM returns to its original vertical position.
The adhesion forces as the AFM tip retracts are minimal comparedto the bulk mechanical loading and, therefore, were not includedin the following analysis [44].
The single particle rupture data was used to verify measurements made on bulk particle films. Fig. 5(a) shows the measured
force per particle, taken as point 3 in Fig. 4(d)–(e), with respect
to particle diameter. Primary sources of error come from the AFM
force data that are subject to the local failure properties of the shelland imperfect centering of the tip contact at the poles of each particle [45]. Due to the differential centrifugation process used toisolate ranges of particle sizes for preparation of the monolayer,
particles of less than 100 nm in diameter or larger than 500 nm
were not tested. However, the AFM data further supports our claimthat the trend is linear by nature and also validates the assumptions of the macroscale model, as seen by comparing the rupture
force magnitude in Figs. 2(b) and 5(a) (combined in Supporting Information Figure S5). We believe that the slight difference of theslopes between the film sintering experiments and the single particle rupture experiments can be attributed to the difference in tip
geometry and thus, the surface area in contact with the plate or tip[45–47]. When rupture occurs during compression between parallel plates, the contact area is much larger than with a sharp tip,where the contact area on the particle is only a small fraction of thetotal surface area [19,45].
3.4. Particle shell characterization
The individual particle rupture experiments were further used
to derive the stiffness of the particle shells. The particles can beconsidered as isotropically-elastic thin-shell microspheres as longas: (1) the ratio of shell thickness to radius is smaller than 0.05,(2) the applied load is point-like, and (3) the load is exerted at
the poles of the particle. The popping experiments completed here
fulfill the former two conditions and partially fulfill the third,
as shown by the centered application of force on the particle inFig. 4(c) [43,48]. Meeting these requirements, the particle shell
stiffness is the slope of the force–distance-curve, defined as thelinear region between points (2) and (3) in Fig. 4(d)–(e):
k=

f

δ

(2)

where k is the stiffness of the particle shell, f is the measured
rupture force for a particle, and δ is the displacement of the AFM
tip. A plot of resulting stiffness for all particles ruptured can be seenin Fig. 5(b). The average shell stiffness of all liquid metal particles

tested was 16.04 ± 4.03 N/m, which is comparable to the stiffnessof harder foams and softer polymers (E ∼ 0.4 GPa). Based on thesefindings, we conclude that there is no correlation between thesize of thiolated particles and stiffness of the thiol shell. However,the derived thiol shell stiffness is useful for both approximating
the stiffness of metal oxide shells for small particle diameters(<100 nm, as shown in Fig. 2(b)), as well as identifying our ability
to rupture the particles on various substrates. Throughout ourexperiments, we have qualitatively noted that it is easier to rupture
the liquid metal nanoparticles on stiffer substrates, while rupture
is often not achievable on very soft substrates. We suspect that
this observation is directly correlated to the stiffness of the particle
shell relative to the stiffness of the substrate. For maximally
flexible electronics printed on soft elastomers, larger, oxide-coatedparticles will minimize the effect of stiffness of the particle shelland enable rupture compatibility with soft substrates.
4. Conclusion
In this study, we have extended previous characterization ofliquid metal nanoparticles based on exposure to ultrasonication,
specifically to achieve particles between 650 and 50 nm indiameter. The electrical response of thiolated and non-thiolatedparticle films was evaluated via mechanical compression between
parallel plates, demonstrating a threshold force at which sintering
may be achieved. A linear relationship between particle size andparticle rupture force was derived and validated via the ruptureof individual particles using AFM. Furthermore, we have derived
the relationships between particle formation, core composition,
shell stiffness, and conductance after rupture and coalescence.
We expect these relationships to contribute to future scalable
manufacturing of liquid metal-based soft electronics.
5. Experimental section*Portions of this procedure, specifically regarding nanoparticle
fabrication and film rupture testing, were originally conducted by
Boley et al. [19] They are shown here as a reference.*
5.1. Liquid metal nanoparticle fabrication
All non-thiolated liquid metal nanoparticles were made using afixed mass (362 ± 5 mg) of eutectic gallium–indium (EGaIn; Sigma
Aldrich PN 495425) deposited into a Kimble Chase 3 dram glass vial
using a 3 mL BD Syringe and a Nordson 14G needle (PN 7018043),
weighed using a OHAUS Pioneer mass balance. Prior to the additionof materials into the glass vial, a sequential cleaning process
consisting of washing and drying Liquinox cleaning solution(Alconox PN 1232), distilled water, and neat ethanol was done toprevent any interference from chemical residue on the interior of
the vial. Each sample was suspended in the vial by micropipetting(BioPette Plus PN BPP1000) 4 mL of ethanol (Koptec PN V101).
The suspended samples were made using a QSonica Q700 microtip
sonicator (PN 4417) at 30% amplitude at a fixed distance (≈1 mm)
for various durations, ranging from 30 to 960 min. Parafilm wasplaced over the opening of the vial to prevent foreign entry. Priorto obtaining samples used for particle or film characterization,
each sample was vigorously re-suspended using a Vortex-Genie 2
shaker (Cole–Parmer PN UX-04724-05) to achieve uniform particle
distribution within the suspension.
For creation of thiolated liquid metal nanoparticles, a diluted
thiol ethanol solution was added to the cleaned glass vials
containing the fixed mass of liquid metal. The solution was createdby adding a measured mass of 3-mercapto-N-nonylpropionamid
thiol (Sigma Aldrich PN 686492) to a volume of ethanol andstirred vigorously until fully dissolved. Samples were sonicatedT.R. Lear et al. / Extreme Mechanics Letters 13 (2017) 126–134

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Fig. 4. Schematics of the AFM experiment single particle rupture experiment. (a) Self-assembled monolayer of liquid metal nanoparticles using the Langmuir–Schaefer
method. Scale bar represents 3 µm. (b) AFM Z map (top), AFM phase map (middle), and SEM images (bottom) of particles before (left) and after (right) rupture. Images on
the right show that the particle shell is visibly broken and liquid metal has spilled out. Scale bars are 100 nm in length. (c) Cross-section of a single liquid metal nanoparticle
before and after particle rupture. (d) Force–displacement curve represents a single popping test. Numbers correspond to process in (e). (e) Schematic showing particle
popping sequence: 1. AFM tip makes contact with particle; 2. Viscoelastic shell of particle undergoes limited deformation; 3. AFM tip penetrates oxide shell, liquid metal is
exposed; 4. Complete fracture of oxide shell, liquid starts to flow out; 5. Exposed liquid reforms shell, causing variation in force as shell resists tip motion.

Fig. 5. (a) Single particle rupture data—measured rupture force vs. particle diameter. Confidence band around data represents one standard deviation based on overall data
set. Linear approximation shown. (b) Stiffness measurements for all AFM particle rupture tests. The dotted line represents the overall average.

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for 60 min at room temperature before undergoing a differential
centrifugation process to focus efforts on a limited range of particle
sizes. The samples were first subjected to 6238 relative centrifugal
force (RCF) in a micro centrifuge (Cole–Parmer PN WU-39065-05)
for 20 min. The aliquot was removed and replaced by an equalvolume of new ethanol, in which the pellet was resuspended using
a bath sonicator (Branson M1800) and the Vortex-Genie shaker.
The samples were then centrifuged again, but at only 250 RCF.
This process was repeated at this RCF three times in order toachieve particle size purification. Thiolated particle shells wereanalyzed using X-ray photoelectron spectroscopy (XPS), showing
the presence of some oxide in the shell composition (see Table S1
in Supporting Information).
5.2. Particle characterization
Scanning electron microscopy (SEM) was used to determine theaverage particle size. After depositing 50 µL of each onto cleaned(rinsed/dried with acetone, isopropanol, ethanol, water) siliconwafers (WRS Materials PN 3P0SSP-002) using a micropipette(BioPette Plus BPP200), high resolution images were obtained using an SEM (Philips XL-40 FEI, 15 kV, 3 µm beam spot size). The
images were then processed using ImageJ (FIJI) through the following procedure: (1) Raw grayscale SEM images were converted to abinary image (Image > Adjust > Threshold, with a lower cutoffof ≈30 and an upper cutoff of ≈230 and dark background option);(2) each resulting image was segmented using built in ImageJ watershed function (Process > Binary > Watershed) [49]; (3) thesizes of all visible particles were calculated using ImageJ’s particle
analysis package (Analyze > Analyze Particles with size ranging
from 10 square pixels to infinity, circularity from 0 to 1, showing
the bare outline of particles); (4) the diameter was computed foreach particle based on a spherical approximation and then averaged for each sample.
5.3. Film rupture—conductance testing
The non-thiolated dispersion was added one drop at a time
(50 µL in total) via micropipette (BioPette Plus BPP200) into areservoir made of polydimethylsiloxane (PDMS) elastomer. After
spin coating PDMS (Dow Corning Sylgard 184) with a 10:1 ratio
between the base and crosslinking agent at 200 RPM using a spincoater (Specialty Coating Systems G3-8) copper wires were added
to each reservoir. A second coat of PDMS was applied using theabove spin coating process. Using a desktop laser system (Universal
Laser Systems VLS2.30), ≈2.5 mm holes were cut using a laser
ablation technique to hold the added dispersion later on (FigureS6). The samples were loaded into a single column materials tester(Instron PN 3345) fitted with a 1 kN load cell. The copper leadson the PDMS reservoirs were connected to a Wheatstone bridge
circuit, where the other three resistors in the circuit were 1000 
and the voltage supply at 5 V (Figure S7). The circuit output wasconnected to the materials testing unit so that the built in software
package could measure the external voltage output as a function ofboth displacement and applied force. Samples were compressed ata steady rate of 1 mm min−1 three times for each desired test.
5.4. Oxide phase validation
Non-conductive samples underwent identical compression
tests using the materials tester but with a heated base. A heat
sheet (McMaster-Carr PN 8009T14) was wrapped around thesteel base of the materials tester and heated the sample via
conduction through the material. The temperature was elevatedusing a programmable temperature controller (McMaster-Carr PN

4314K6) and verified with an attached thermocouple (McMasterCarr PN 37045K192). The temperature of the materials tester base
was held constant at 50 °C. After depositing 50 µL of dispersion into
the reservoir, the compression test resumed as described above at
a compression rate of 1 mm min−1 .
5.5. Langmuir monolayers of liquid metal nanoparticles
A thiolated dispersion of liquid metal nanoparticles was sonicated for one hour before centrifugation and added to a Kibron
microtrough (5 cm × 23 cm). As seen in Figure S4, particle size iscontrolled by differential centrifugation, a process that uses various g-forces to fractionate the particles. By using high forces andresuspending the precipitate, the unwanted under-sized particles
can be discarded. Likewise, exposure to lower g-forces and keeping the supernatant allows for the disposal of over-size particles.
This method allows us to create a monolayers of particles withminimal size variations. For each monolayer, a total of 3 µL of thiolated dispersion (concentration of 50 mg/mL) was added to thesubphase one drop at a time—approximately 0.6 µL/s using a micropipette (BioPette Plus BPP200). The temperature of the deionized water (18 M cm) in the trough was 20 °C. Once the ethanol
carrier solvent evaporated, the hydrophobic thiolated liquid metal
nanoparticles remained on the subphase and a reflective dark silver was observed (reflectivity points to smooth surfaces). The sample was then compressed at 10 mm/min to improve surface coverage (Figure S8). Using the Langmuir–Schaefer method, the newly
formed monolayer films were transferred onto an OTS-modified
silicon wafer substrate. In order to characterize both the particle
size and polydispersity of the nanoparticles and the films composed of them, dynamic light scattering (DLS) was employed using a Malvern Zetasizer Nano ZS at a scattering angle of 173°. SEM
imaging was performed on the films using a NOVA nanoSEM FESEM field-emission with an accelerating voltage of 10k. The image
characterization process was identical to that of the Particle Characterization section. AFM height maps were taken of each film using specifications found in Table S2.
5.6. Single particle rupture
Langmuir–Schaefer monolayers were deposited onto an OTS
modified silicon wafer substrate (without pt coating) as previously
described. The temperature of the deionized water (18 M cm)
in the trough was 20 °C. SEM observations were performed using
a NOVA nanoSEM FESEM field emission (FEI) high-resolution microscope using an accelerating voltage of 10 kV. AFM topographic
images and force–distance curves were acquired on a Multimodescanning probe microscope (Veeco, USA) at ambient conditions.
The sharp indentation tip used was a hard-diamond-like-carbon(Figure S9). The spring constants of the probe were determined using the thermal tune method [50]. The force measurements weredone using probe HQ NSC 15/Hard/AL BS (MikroMasch) cantileverwith a force constant of 40 N/m, n-type silicon tip (h = 12–18 µm,
r < 20 nm) (Table S2). The substrate surface was reanalyzed after the tip had both come into contact and been retracted (Figures S10–S11). For the force–distance curve analysis, the sensitivity and spring constant of the cantilever were kept constantat 41.95 ± 2.96 nm/V and 30.40 ± 0.44 N/m, respectively. Procedures were reviewed for: force–displacement curve measurement [50–52], properties of nano-sized capsules [53], adhesion
force curve analysis [54], Atomic J software [55], AFM mechanicalproperties [56], AFM colloid properties [57], and measurements ofrigidity [58].
Acknowledgments
We would like to thank Michelle Yuen for her willingness tohelp with MATLAB coding and experiment design, Adam Bilodeau

T.R. Lear et al. / Extreme Mechanics Letters 13 (2017) 126–134

for his help with SEM, and Mohammed Mohammed for his
suggestions regarding various particle characterization methods
and figure preparation.
This work was primarily support by a National Science
Foundation CAREER Award (1454284). J.W.B. was supported by aNASA Space Technology Research Grant (NNX14A052G). E.W. was
supported by a National Science Foundation Graduate Research
Fellowship (DGE-1333468).
Author contributions
T.L. conducted the film experiments and discussed results with
E.W. S.H. and J.W.B. conducted experiments with the Langmuir
trough and AFM. T.L. wrote the manuscript with help from S.H. and
R.K. Further discussions and theories were developed with D.T. and
R.K.
Appendix A. Supplementary data
Supplementary material related to this article can be foundonline at http://dx.doi.org/10.1016/j.eml.2017.02.009.
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