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Ultralight, ultrastiff mechanical metamaterials
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Zheng, Xiaoyu, Howon Lee, Todd H. Weisgraber, Maxim
Shusteff, Joshua DeOtte, Eric B. Duoss, Joshua D. Kuntz,
Monika M. Biener, Qi Ge, Julie A. Jackson, Sergei O. Kucheyev,
Nicholas X. Fang, and Christopher M. Spadaccini. "Ultralight,
ultrastiff mechanical metamaterials." Science (20 June 2014)
344:6190 pp. 1373-1377.
As Published
American Association for the Advancement of Science (AAAS)
Author's final manuscript
Wed Sep 03 15:25:11 EDT 2014
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Submitted Manuscript: Confidential
07 May 2014
Title: Ultra-light, Ultra-stiff Mechanical Metamaterials
Authors: X. Zheng1*, H. Lee2, T. H.Weisgraber1, M. Shusteff1, J. Deotte1, E. B. Duoss1, J. D.
Kuntz1, M. M. Biener1, Q. Ge2, J. A. Jackson1, S. O. Kucheyev1, N. X. Fang2*, C. M.
Lawrence Livermore National Laboratory, Livermore, CA 94550, USA.
Department of Mechanical Engineering, Massachusetts Institute of Technology, Cambridge,
MA 02139, USA.
*Correspondence to: X. Z.: [email protected], C.M.S.: [email protected], N.X.F.:
[email protected]
Abstract: The mechanical properties of ordinary materials degrade substantially with
reduced density, due to the bending of their structural elements under applied load. We
report a class of micro-architected materials that maintain a nearly constant stiffness per
unit mass density, even at ultra-low density. This performance derives from a network of
nearly isotropic microscale unit cells with high structural connectivity and nanoscale
features, whose structural members are designed to carry loads in tension or compression.
Production of these microlattices, with constituent materials ranging from polymers to
metals and ceramics, is made possible by using projection microstereolithography, an
additive micromanufacturing technique, combined with nanoscale coating and postprocessing. We found that these materials exhibit ultra-stiff properties across more than
three orders of magnitude in density, regardless of the constituent material.
One Sentence Summary:
We report a class of micro-architected materials that change their stiffness linearly with reduced
Main Text: Nature has found a way to achieve mechanically efficient materials by evolving
cellular structures. Natural cellular materials, including honeycomb (1) (wood, cork) and foamlike structures, such as trabecular bone (2), plant parenchyma (3), and sponge (4), combine low
weight with superior mechanical properties. For example, lightweight balsa has a stiffness-toweight ratio comparable to that of steel along the axial loading direction (5). Inspired by these
naturally occurring cellular structures, manmade lightweight cellular materials fabricated from a
wide array of solid constituents are desirable for a broad range of applications including
structural components (6, 7), energy absorption (8, 9), heat exchange (10, 11), catalyst supports
(12), filtration (13, 14), and biomaterials (15, 16).
However, the degradation in mechanical properties can be drastic as density decreases (17, 18).
A number of examples among recently reported low-density materials include graphene
elastomers (19), metallic micro-lattices (20), carbon nanotube foams (21), and silica aerogels
(22, 23). For instance, the Young’s modulus of low-density silica aerogels (22, 23) decreases to
10 kPa (10–5 % of bulk ) at a density of less than 10 mg/cm3 (< 0.5% of bulk).
This loss of mechanical performance is because most natural and engineered cellular solids with
random porosity, particularly at relative densities less than 0.1%, exhibit a quadratic or stronger
scaling relationship between Young’s modulus and density as well as between strength and
density. Namely, E/Es ~ (/s)n and yys ~ (/s)n, where E is Young’s modulus,  is density, y
is yield strength, and s denotes the respective bulk value of the solid constituent material
property. The power n of the scaling relationship between relative material density and the
relative mechanical property depends on the material’s microarchitecture. Conventional cellular
foam materials with stochastic porosity are known to deform predominantly through bending of
their cell walls and struts (24). This type of deformation results in relative stiffness scaling with n
= 2 or 3. A number of approaches in recent years have aimed to reduce this coupling between
mechanical properties and mass density (5, 17, 18, 20, 25-31). Among these, few fabrication
processes are capable of building arbitrary three-dimensional microarchitectures with controlled
micro- and nanostructure across a wide range of mass density and material constituents. The
desired material properties are thus limited to a narrow density range and specific loading
Improved mechanical properties can arise from a material that contains micro- and nano-scale
building blocks arranged in an ordered hierarchy. Among these new designs are metallic
microlattices with high recoverability when compressed (20, 26), TiN nanotrusses (32, 33) and
ceramic composite trusses (34) that show enhanced fracture toughness of coating materials when
the thickness of coating materials is reduced to the nanoscale.
We report a group of ultra-light mechanical metamaterials that maintain a nearly linear scaling
between stiffness and density over three orders of magnitude in density, over a variety of
constituent materials. We use the term “mechanical metamaterials” to refer to materials with
certain mechanical properties defined by their geometry rather than their composition. The
materials described here are highly ordered, nearly isotropic, and have high structural
connectivity within stretch-dominated, face-centered cubic architectures. The ultra-low density
regime is accessed by fabricating microlattices with critical features ranging from ~20 m down
to approximately ~40 nm. The densities of samples produced in this work ranged from 0.87
kg/m3 to 468 kg/m3, corresponding to 0.025–20% relative density.
A stretch-dominated unit cell structure, consisting of b struts and j frictionless joints and
satisfying Maxwell’s criteria, M= b – 3j + 6 > 0, is significantly more mechanically efficient,
with a higher stiffness-to-weight ratio (defined as E/) than its bend-dominated counterpart. This
is attributed to its struts carrying load under compression or tension rather than bending(17). A
fundamental lattice building block of this type is the octet-truss unit cell (Fig 1 (A)) whose
geometric configuration was proposed by Deshpande et al.(35). The cell has a regular
octahedron as its core, surrounded by eight regular tetrahedra distributed on its faces (Fig. S1).
All the strut elements have identical aspect ratios, with twelve connected at each node. The cell
has a face centered cubic (FCC) structure whose cubic symmetry generates a material with
nearly isotropic behavior (36). The relative density of such octet-truss unit cells can be
approximated by = 26.64 (d/L)2 (35), where L and d are the length and diameter of each beam
element. On the macroscale, under uniaxial compressive loading, the relative compressive
stiffness and yield strength of these structures theoretically show linear scaling relationships:
E/Es ~/s) and /s ~/s) (35). A cubic lattice is readily constructed by periodic packing of
the unit cell along its three principal directions (Fig. 1(B) and (C)) (37, 38). Alternate
orientations of the bulk lattice relative to the unit cell’s principal axes can likewise be
constructed (Fig. S2), with the fundamental tessellation of space by the unit cell remaining the
same. To study how the loading direction and lattice orientation affect their E~ scaling
relationship, octet-truss lattices were analyzed, fabricated, and tested in a variety of orientations
((39), Fig S1-Fig S5). In addition to these stretch-dominated lattices, as a point of comparison, a
bend-dominated tetrakaidekahedron unit cell (40, 41) of the same size scale was generated and
the corresponding cubic-symmetric foams (known as Kelvin foams), were fabricated with a
variety of densities (Fig. 1(D–F)).
The fabrication of these microlattices is enabled by projection microstereolithography, a layerby-layer additive micromanufacturing process capable of fabricating arbitrary three-dimensional
micro-scale structures (42, 43). In contrast to other 3D rapid prototyping methods such as 3D
printing and UV projection waveguide systems (44), this type of fabrication technology is ideal
for 3D lattices with high structural complexity, and with feature sizes ranging from tens of
microns to centimeters. By combining projection microstereolithography with nano-scale coating
methods, 3D lattices with ultra-low relative densities below 0.1% can be created. The process
begins with a photosensitive polymer resin bath; we use either 1, 6-hexanediol diacrylate
(HDDA) or poly (ethylene glycol) diacrylate (PEGDA). Shown schematically in Fig. 2(A), the
apparatus uses a spatial light modulator—in this case a liquid-crystal-on-silicon chip—as a
dynamically reconfigurable digital photomask. A three-dimensional CAD model is first sliced
into a series of closely spaced horizontal planes. These two-dimensional image slices are
sequentially transmitted to the reflective liquid-crystal-on-silicon chip, which is illuminated with
UV light from a light emitting diode array. Each image is projected through a reduction lens onto
the surface of the photosensitive resin. The exposed liquid cures, forming a layer in the shape of
the two-dimensional image, and the substrate on which it rests is lowered, reflowing a thin film
of liquid over the cured layer. The image projection is then repeated with the next image slice
forming the subsequent layer. Our polymer microlattices were fabricated in tens of minutes and
have features spanning size scales from 10 to 500 m. For mechanical testing purposes, all
materials described here were fabricated as blocks of various sizes consisting of multiple unit
cells (Table S1). Fig. 2(B) and (F) show scanning electron micrographs (SEM) of the as-built
polymer lattice and unit cell.
Although projection microstereolithography requires a photopolymer, other constituent
materials, such as metals and ceramics, can be incorporated with additional processing. Using the
base polymer lattice as a template, we are able to convert the structures to metallic and ceramic
microlattices. Metallic lattices were generated via electroless nickel plating on the as-formed
HDDA. The thickness of the metal coating is controlled by the plating time, yielding metal films
from 100 nm to 2 m. The polymer template is subsequently removed by thermal
decomposition, leaving behind the hollow-tube nickel-phosphorus (Ni-P) stretch-dominated
microlattice shown in Fig. 2(C) and (G).
A similar templating approach is used to generate hollow-tube aluminum oxide (amorphous
Al2O3, alumina) microlattices; however, the coating is produced by atomic layer deposition
(ALD), a gas phase process, rather than liquid phase processing. The resulting hollow-tube
microlattices have alumina thicknesses from ~40 to 210 nm, with an example shown in Fig. 2(D)
and (H), with corresponding material weight density ranging from less than 1 kg/m3 to 10.2
Loading the resin bath with nanoparticles can further expand the base material set. Solid Al2O3
ceramic lattices were prepared in the microlithography system by using photosensitive PEGDA
liquid prepolymer loaded with ~150 nm alumina nanoparticles (Baikowski Inc, approximately
12.5% alumina by volume). The same sequential lithographic exposure process produced a
microlattice made of a hybrid of solid PEGDA and alumina nanoparticles. These hybrid lattices
are converted to pure Al2O3 octet-truss microlattices through a sintering procedure (39). An
example of this structure is shown in Fig. 2(E) and (I). The parameters and properties for a
selection of our stretch-dominated mechanical metamaterials and bend-dominated foams are
summarized in Table S1. The densities of all samples were calculated by measuring the weight
and fabricated dimensions of the completed microlattices.
The microstructured mechanical metamaterials were tested to determine their Young’s modulus
E and uniaxial compressive strength y defined as the crushing stress of the material.Uniaxial
compression studies of all microlattices with the same cubic dimensions were conducted on an
MTS Nano Indenter XP, equipped with a flat punch stainless steel tip with a diameter of 1.52
mm. During 20 consecutive compression cycles up to 10% strain, we observed typical
viscoelastic behavior for the polymer microlattices with pronounced hysteresis with loading rate
dependent Young’s modulus. The Young’s moduli for all polymer microlattices and foams were
extracted at loading rate at 87.2 nN/s, corresponding to a strain rate of 10-3 s-1. Uniaxial
compression of these structures is shown in Movies S1-S3. Representative stress-strain curves
from uniaxial compression crushing tests for determining the compressive strength of octet-truss
microlattices made of solid HDDA polymer, hollow-tube Ni-P metal, and solid alumina are
shown in Fig. S7 (A), (B) and Fig. S8 (A) respectively. Bulk HDDA polymer, cured by UV
crosslinking a solid sample of similar dimensions to the octet-truss lattices, was determined to
have Young’s modulus and yield stress values of Es = 530 MPa and σys= 86 MPa, respectively.
The bulk property values for other constituent materials and the detailed methods of
measurement are given in supplementary sections (39).
The results of these mechanical tests, together with the bend-dominated tetrakaidekahedronbased Kelvin foams fabricated from the base HDDA polymer, are summarized in Fig. 3(A) and
Fig. 3(B), which plot the relative Young’s modulus and strength against the relative mass
density, respectively. Fig. 4 shows the location of these material properties on the stiffness
versus density material selection chart, together with other recently reported ultra-light materials
for comparison. The stretch-dominated microlattices populate the highly desirable ultra-light,
ultra-stiff space toward the upper left of the chart (17) and have stiffness-to-weight ratios that do
not substantially degrade as density decreases by several orders of magnitude. In contrast to the
common bend-dominated E/Es~(s)2 scaling of open-cell stochastic foams such as silica
aerogels and carbon foams, our stretch-dominated microlattice materials demonstrate the desired
linear relationship of E/Es~s, approaching the theoretical limit, and exhibit this remarkable
scaling relationship over three orders of magnitude in density and across all constituent materials
studied. These octet-truss lattice materials are highly isotropic, so the scaling of stiffness with
density does not vary with the orientation of the lattice (Fig S4), as confirmed by our studies of
different loading directions. These lattices have the highest specific stiffness when the lattice is
loaded normal to the (111) plane, which is closest-packed within the FCC architecture. 
In the ultralow-density regime (relative density < 0.1%), we observed markedly different
compression behavior in hollow-tube ALD ceramic octet-truss microlattices, compared to solid
ceramic lattices at higher relative densities (8–20%). The hollow-tube ceramic microlattices with
nanoscale wall thicknesses showed smoother behavior with progressively fewer discontinuities
in their stress-strain curves (Fig. S8 (A) and (B)) in contrast to solid microstrut ceramic lattices
with catastrophic, fracture-dominated behavior. The loading-unloading curves of hollow-tube
Al2O3 lattices revealed elastic behavior followed by a nonlinear response on each loading cycle.
Although relative compressive stiffness and relative density initially follow a nearly linear
scaling law, the transition from conventional brittle behavior (in low-density ceramic materials)
to more “ductile” mechanical behavior (in ultra-light materials with nanoscale wall thicknesses)
suggests a transition from a fracture-dominated failure mode to a buckling-dominated failure
mode with suppression of the catastrophic failure seen in solid Al2O3 octet-truss lattices. These
differences in compressive behavior between solid and hollow-tube ceramic octet-truss lattices
are primarily attributed to local buckling induced by the high aspect ratio of the strut length to
nanoscale wall thickness, in contrast to nanoscale TiN trusses (32) and ceramic composite (34)
where the aspect ratio is low enough to allow the nanoscale strengthening effect of the wall
thickness to dominate. For example, the ratio of strut length to nanoscale wall thickness in Fig.
S8(B) is approximately 1400:1 and contributes to its large compression strain, governed by
Thus, the relative compressive strength transitions from the nearly linear scaling
law governing the stretch-dominated failure mode at an approximate density near 0.08% to a
scaling power of 2.7, as indicated in Fig. 3(B). In the same figure, a similar transition from
yielding dominated to buckling dominated failure at an approximate relative density of 0.2% is
evident in Ni-P lattices, consistent with the trend observed on bend-dominated metallic
microlattices (26).
When an ultralow-density metallic microlattice is bend-dominated, its stiffness degrades
significantly with reduced density. An example of this is the Ni-P lattice reported by Schaedler
et al. (20), whose specific stiffness (stiffness to weight ratio) degrades from 0.23 x 106 m2/s2 to
0.05 x 106 m2/s2 as density is reduced from 40 mg/cm3 to 14 mg/cm3 (45). Significantly, our Ni-P
stretch-dominated metallic lattice is not only much stiffer in the same density range, its specific
stiffness stays nearly constant, measured as 1.8 x 106 m2/s2 and 2.1 x 106 m2/s2 at densities of 14
mg/cm3 and 40 mg/cm3, respectively. Similarly, in a recent report of high-strength
microarchitected ceramic composites (34), their strength performance approaches the linear
scaling relationship over a narrow density range, and only when loaded in a direction optimized
for their anisotropic architecture. Our metamaterials, in contrast, maintain their mechanical
efficiency over a broad density regime, without substantial degradation in specific stiffness,
owing to the nearly linear E~ scaling relationship.
We have shown that these high mechanical efficiencies are possible across a range of constituent
materials. Fabricating ordered lattice structures at these length scales brings them into the regime
in which it becomes possible to interact with the structure as a material.
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Acknowledgments: This work performed under the auspices of the U.S. Department of Energy
by Lawrence Livermore National Laboratory under Contract DE-AC52-07NA27344.
Funding support from LDRD Strategic Initiative 11-SI-005 and DARPA MCMA
(Materials with Controlled Microstructural Architecture, Program Manager Dr. Judah
Goldwasser) is gratefully acknowledged. The authors at Lawrence Livermore National
Laboratory wish to acknowledge Howard Rathbun for useful discussions on modeling the
lattice structures as well as Dr. Marcus Worsley for taking SEM images and Christopher
Harvey for his technical support (LLNL-JRNL-640334).
Fig. 1. Architecture of stretch-dominated and bend-dominated unit cells and lattices. (A)
Mechanical response to compressive loading of a stretch-dominated octet-truss unit cell. (B)
Octet-truss unit cells packed into a cubic microlattice. (C) SEM image of a stretch-dominated
lattice material that is composed of a network of octet-truss unit cells. (D) Mechanical response
to compressive loading of a bend-dominated tetrakaidekahedron unit cell. (E)
Tetrakaidekahedron unit cell packed into a cubic bend-dominated lattice (Kelvin foam). (F) SEM
image of a bend-dominated lattice that is composed of a network of tetrakaidekahedron unit
Fig. 2. Fabrication of ultra-light, ultra-high stiffness stretch-dominated microlattices. (A)
Projection microstereolithography, a layer-by-layer technique capable of fabricating arbitrary,
microscale, three-dimensional structures with ~5 m resolution; the SEM imageon the right is
the fabricated octet-truss unit cell. (B)–(E) Octet-truss microlattices with varied constituent
materials and configurations: (B) solid polymer HDDA. (C) hollow-tube metallic Ni-P, (D)
hollow-tube ceramic (alumina) (E) solid ceramic (alumina). (F)-(I) are magnified views of the
struts of the microlattices in (B)-(E), respectively.
Fig. 3. (A) Relative stiffness as a function of relative density for stretch-dominated microlattices
and bend-dominated microlattices. Stretch-dominated materials exhibit a linear stiffness–density
relationship of E~, whereas bend-dominated materials soften as E~2 or worse, as shown by
their slopes in the plot. (B) Relative strength as a function of relative density for stretchdominated and bend-dominated microlattices.
Fig. 4. An Ashby chart plotting the compressive stiffness vs. the density of ultra-light, ultra-stiff
mechanical metamaterials compared to other previously reported materials. Dotted lines indicate
contours of constant stiffness–density ratio c = E/(m2/s2).
Supplementary Materials:
Materials and Methods
Figures S1-S9
Table S1
Movies S1-S3
References (46-55)
Supplementary Materials for
Ultra-light, Ultra-stiff Mechanical Metamaterials
Xiaoyu Zheng1*, Howon Lee2, Todd H Weisgraber1, Maxim Shusteff1, Joshua R. Deotte1, Eric
Duoss1, Joshua D. Kuntz1, Monika M. Biener1, Qi Ge2, Julie A. Jackson1, Sergei O. Kucheyev1,
Nicholas X. Fang2*, Christopher M. Spadaccini1*
Correspondence to: Christopher M. Spadaccini ([email protected]),
Nicholas X. Fang ([email protected])
Xiaoyu Zheng ([email protected])
This PDF file includes:
Materials and Methods
Figs. S1 to S9
Table S1
Captions for Movies S1 to S3
Other Supplementary Materials for this manuscript includes the following:
Movies S1 to S3
Materials and Methods
Octet-truss lattice geometric configurations
The octet-truss lattice structures in this study satisfy the Maxwell deterministic stability
criteria for a stretch-dominated pin-jointed frame (17, 46). As described in the main article, the
face-centered cubic (FCC) unit cell has a regular octahedron core with regular tetrahedra affixed
to each face, such that twelve strut elements connect at each node. The unit cell’s cubic
symmetry (detailed in Figure S1) results in a material with nearly isotropic properties (36). The
microstructured lattice material is produced by “stacking” unit cells with each strut shared
between two neighboring cells – i.e. periodically repeating unit cell elements along three
orthogonal axes. We refer to different lattice orientations by the Miller indices of the lattice plane
that forms a flat face to which loading is applied. Thus, a (111)-oriented lattice has a (111) plane
at its top surface, to which uniaxial compressive loading is applied in the [111] direction.
We carried out finite-element (FEM) studies of four different lattice orientations: (001) ,
(111) , (112) , and (110) shown in Fig.S2. These studies complemented and expanded upon the
experimental measurements of (001)- and (111)-oriented solid polymer HDDA lattices,
fabricated by projection micro-stereolithography. The finite element simulations were
performed with the NIKE3D fully implicit solver, developed at Lawrence Livermore National
Laboratory (47). Each strut in the octet-truss lattice was represented by a single Hughes-Liu
beam element (Fig S3). The simulations for the [001] loading direction contained 27 repeating
unit cells and 172 nodes. For the other loading directions, the FEM model was based on the
geometry of the stereolithography-fabricated part for the (111) orientation (Figure S2 compare
(B), (C) and (D)) and contained 265 nodes. Using a linear elastic materials model based on the
properties of HDDA, the relative stiffness at each density was derived from the slope of the
stress response with applied loading. To mimic the experimental conditions, for each loading
direction, boundary conditions were established for two sets of nodes contained within defined
upper and lower planar surfaces with normals coincident with the load (arrows in Fig. S3). A
linear displacement was applied to the nodes of the upper plane, while the lower plane nodes
remained fixed. The remaining nodes were unconstrained.
It is clear from the FEM results in Fig. S4 that the overall linear scaling relationship does
not change as the major orientation changes. This confirms that the overall stretch-dominated
behavior operates regardless of loading direction. However, the linear proportionality constant
degrades as the direction deviates from (111) plane which offers the highest stiffness-to-weight
 C( hkl )
ratio. These scaling stiffness constants in different orientation planes, namely,
found to be C( 001) = 0.2412, C(111) = 0.3884, C(110 ) = 0.2269, and C(112 ) = 0.1596.
These findings are supported by mechanical testing of real structures in a subset of these
orientations (Fig. S5(A)). The results of uniaxial compression tests on the solid polymer HDDA
lattices with (001) and (111) orientations, carried out as detailed in the main article, are
summarized in Fig. S5(B), and compared with the directional sensitivity of honeycomb
structures and bend-dominated open-cell foam structures. It is apparent that the honeycomb,
when loaded out of plane – i.e. perpendicular to the axis of the hexagonal prisms (29) has the
optimal linear scaling between E and r, whereas when loaded in-plane, its stiffness suffers
greatly (25, 29). This was confirmed in a recent study of composite polymer-ceramic
microarchitectures, where only a honeycomb optimized for out-of-plane loading exhibited the
ideal linear stiffness-density scaling relationship (34). In contrast, the octet-truss material, due to
its symmetry, exhibits optimal linear scaling of E with r for all of the orientations studied here,
with the (111) orientation slightly stiffer than the (001) orientation.
Octet-truss lattice defect sensitivity
Defects in the lattice material building blocks may have crucial effects on the macroscopic
properties of the material, and there are numerous prior studies of other cellular materials. For
honeycomb structures, for example, Ruffoni et al. thoroughly investigated the relationship
between geometric defects and their macroscopic properties, and calculated the strain energy
density around a single defect within two-dimensional cellular solids (48). This has important
implications for the honeycomb, which is one of the most widely used cellular structures, and
sheds light on other 2D and 3D cellular structures. Symons and Fleck (49) present a detailed
treatment of the role of various microstructural imperfections (missing struts, misplaced joints,
and deformed cell walls) upon the in-plane effective properties of three isotropic lattices:
hexagonal, triangular and Kagome. Alkhader and Vural compare the influence of defects in
honeycomb structures, with tetragonal and triangular lattices (50). The weakening of mechanical
properties due to misplaced joints is numerically explored in a parallel study by Romijn and
Fleck for three isotropic lattices, as well as two orthotropic topologies: square and diamondcelled (51). All these studies conclude that the defect sensitivity of modulus and fracture
toughness for each lattice is highly dependent on its nodal connectivity, consistent with the
arguments of Deshpande et al (46). Lattices with a high connectivity value of 6 or more, such as
the triangular lattice, are highly insensitive to defects, compared to square lattices. The
misplaced joints have a negligible effect upon their mechanical properties.
This analysis is highly favorable for the octet-truss lattice, with its exceptionally high nodal
connectivity of 12. In contrast, some lattices, such as the Kagome, have a connectivity value that
is in a transitional range. The response of these lattices can be bending- or stretching-dominated,
depending on the degree of imperfection. For instance, under uniform loading a square lattice
deforms by strut stretching, but upon the introduction of a defect such as a macroscopic crack the
struts deform by a combination of stretching and bending. Thus, the moduli and strength of this
type of topology are highly sensitive to imperfections.
For generality, we chose to study the defect sensitivity of the octet-truss lattices using an
FEM approach. While it is possible to generate parts with controlled defects, this would
significantly limit the number and type of defects within the study. Moreover, misplaced joints
and node positions are the most common defects in 3D manufactured trusses. Here, we model
defects that are displacements of the node positions from their proper locations within the octettruss geometry. We compare a perfect octet-truss lattice with lattices containing joint
mismatches and irregularities in the shape of the trusses, which arise from the displaced node
positions. In order to simulate geometric imperfections from a more generic perspective, rather
than single nodal defects, the positions of all nodes were translated in all three directions by a
distance described by a Gaussian distribution with zero mean and a standard deviation of =520% of the strut length. Then, as in the study of different orientations described above, a linear
displacement was applied to the upper plane nodes, while the lower plane nodes remained fixed,
and all other nodes were unconstrained. For each perturbation amplitude, the sample size was
100 simulations. Fig S6(A) shows models constructed in NIKE 3D with prescribed defect levels
from 0 to 20% of the strut length. Fig S 6 (B) shows the effect of geometric imperfections on the
relative stiffness of the octet-truss lattice in the (001) orientation, with each cluster of points at a
specific relative density spanning the range of defect levels. The sensitivity to defects can be
assessed by examining the slope of the stiffness as a function of defect level (Fig S6(C)). It is
evident from these results that the octet-truss lattice is highly insensitive to geometric
imperfections, thanks to its high nodal connectivity. Figures S6 (B) and (D) confirm that the
octet-truss lattice remains stretch-dominated, with only a slight reduction in its proportional
stiffness constant, even when perturbed at a defect level of 20%. We find no evidence that any
bending mechanism can be activated by introducing defects.
S3. Fabrication of hollow-tube Ni-P microlattices and measurement of bulk properties
All chemicals for electroless nickel deposition were obtained from Sigma-Aldrich
(Milwaukee, WI) unless otherwise indicated. To prepare each sample for electroless nickel
deposition, the sample was first cleaned by immersing in an acetone bath with gentle agitation
for 2 minutes. After rinsing in deionized (DI) water for 30 seconds, the sample was immersed in
a 1 M NaOH solution for 2 minutes followed by rinsing in DI water for 1 minute. Next, the
sample surface was sensitized by immersing in an aqueous solution containing 3.5 g/L of
SnCl2·2H2O and 5.0 mL/L 37% HCl for 2 minutes followed by rinsing in DI water for 30
seconds. After sensitization, the sample surface was activated with palladium catalyst by
immersing in an aqueous solution containing 0.5 g/L PdCl and 4.0 mL/L 37% HCl for 2 minutes
followed by rinsing in DI water for 30 seconds. After activation, the surface was metallized with
nickel by immersing in a bath consisting of equal parts by volume of EdgeMet Electroless Nickel
Kit A and B (Buehler, Lake Bluff, IL). This step was performed at 80°C with vigorous stirring.
The coating thickness was controlled by varying the time of the metallization step between 1–20
minutes. Following this step, the sample was removed and rinsed with DI water for 30 seconds
and allowed to air dry. Subsequently, the polymer interior was removed via thermal
decomposition by heat-treating in a box furnace in air at 500°C for 4 hours. This last step yielded
hollow-tube nickel-phosphorus microlattices.
The mechanical properties for the base electroless plated Ni-P material were measured via
nanoindentation with a Berkovich diamond tip. Indents were carried out on a Ni-P film
approximately 2 µm thick, deposited on a silicon substrate by the same method described above.
The film was subsequently annealed at 500 °C for 4 hours, to mimic the conditions experienced
during preparation of hollow-tube Ni-P microlattices. Averaged values were collected from a
least n=6 different indents on each sample, made at a constant loading strain rate of 0.05 s−1. To
avoid substrate compliance effects, indenter penetration depth was limited to no more than 10%
of the total film thickness. Stiffness (Young’s modulus) and hardness were calculated based on
continuous stiffness measurement (CSM) (52) according to the Oliver-Pharr method (53). The
measured Young’s modulus value of 97 GPa after annealing is toward the low end, but well
within the range of reported values for Ni-P (54, 55). The measured hardness value H=5.2 GPa
was used to calculate the yield strength of Ni-P as σs≈ H/3 = 1.73GPa.
S4. Fabrication of AL2O3 hollow-tube microlattices and measurement of bulk properties
Nanoscale Al2O3 films were deposited by ALD using the well-established trimethylaluminum (AlMe3/ H2O) process in a warm wall reactor (wall and stage temperature of 125ºC).
Long pulse, pump, and purge times (90 seconds each) were used to ensure uniform coatings
throughout the porous material. The film thickness was controlled by adjusting the number of
ALD cycles, and the growth rate per cycle was calculated from the measured mass gain, the
known surface area of lattice structures, and density of ALD Al2O3. This rate was found to be
∼0.25 nm per cycle. The bimaterial hybrid polymer-Al2O3 microlattices were then heat-treated in
a box furnace in air at 500°C for 4 hours to remove the polymer template, leaving behind the
Al2O3 hollow-tube octet-truss lattices.
For characterization of the ALD material Al2O3 hollow-tube octet truss lattices,
approximately 200 nm of Al2O3 was deposited onto a silicon substrate at 125 ºC following the
procedure above. ALD Al2O3 film properties were characterized via nanoindentation similarly to
the methods used for Ni-P described in S3. The values of Young’s modulus and yield strength
measured for ALD alumina are Es = 153 GPa and σs = 2.2 GPa, respectively, with a calculated
mass density of 2900 kg/m3.
S5. Fabrication of solid AL2O3 ceramic microlattices
Ceramic loaded polymer lattices were initially created with 150 nm Al2O3 nanoparticles
(Baikowski, Inc.) loaded at approximately 12.5% alumina by volume in the PEGDA prepolymer.
The ceramic-loaded polymer lattice was fabricated through projection microstereolithography
using the same procedure as described in the main text. Subsequently a sintering process was
used to remove the polymer and densify the alumina. This began with staged calcination in air
(120ºC for water removal, 220ºC for volatile hydrocarbon, 420ºC for combustion of
hydrocarbon, and 600ºC for burning out residual carbon) and a final dwell at 1000ºC for 2 hours.
Then, the sintering process continued to 1300ºC for 2 hours in air with a heating rate of 5 degrees
per minute and cooling rate of 10 degrees per minute. The final product was Al2O3 solid lattices
with approximately isotropic linear shrinkage of 32% in all three dimensions and ligament
shrinkage of 32%. This results in a final sintered material density of 85% (relative to fully-dense
bulk ceramic) for the base material of which the truss is made up. Bulk mechanical properties of
alumina were measured by compression loading of commercially-obtained 99.5% pure alumina
(AD-995, CoorsTek, Golden, CO). The values of Young’s modulus and yield strength from
these measurements are Es = 180 GPa and σs= 1.6 GPa.
Fig. S1.
Cubic symmetry of the octet-truss unit cell. (A) Octet-truss unit cell with an octahedral core, and
each surrounding quadrant containing one tetrahedral subunit. (B) The octahedral core of the
unit cell with eight symmetry-equivalent faces, corresponding to the eight equivalent orientations
in the (111) family. (C) Symmetry of the tetrahedral subunits, and the (111) family of
symmetry-equivalent planes.
Fig. S2
Illustration of the different loading orientations of octet-truss lattices studied in this work.
Lattice orientations are referred to by the Miller indices of the lattice plane to which loading is
applied. A (111)-oriented lattice has a (111) plane at its top surface, to which uniaxial
compressive loading is applied in the [111] direction.
Fig. S3
Illustration of the NIKE3D FEM simulation used to study the directional stiffness dependence of
the octet-truss lattice, indicating the boundary conditions for the simulations, and the calculated
axial forces for each strut. A single unit cell is shown at left, and a 3x3x3 lattice at right.
Fig. S4
The results of FEM investigation showing relative compressive stiffness as a function of relative
density of octet-truss lattices in different loading orientations.
Fig. S5
(A) Two orientations of 1 mm x 1 mm x 1 mm octet-truss lattices fabricated by projection microstereolithography. (B) Relative compressive stiffness of octet-truss lattices in (001) and (111)
orientations, shown as a function of relative density, in comparison with honeycomb structures
loaded both in-plane and out-of-plane, as well as a bend-dominated open cell tetrakaidekahedral
(Kelvin) foam, also fabricated by projection micro-stereolithography. The lines are linear fits to
the measured data points.
Fig. S6
Defect sensitivity study of octet-truss lattices. (A) The nodes of a (001)-oriented lattice are
displaced by a prescribed Gaussian distribution with zero mean and a standard deviation of 520% of the strut length. (B) The straight line is the E~ relationship of an octet-truss lattice
without defects. Each cluster of data points for a given relative density spans the range of
geometric defects from 5% to 20%. (C) A typical detailed view of the relative stiffness as a
function of defect level at one relative density (11%) indicated by the box and arrow in (B). (D)
Relative yield strength vs. relative density of octet-truss lattices with varying levels of geometric
Fig. S7
Representative uniaxial compressive loading and unloading response of octet-truss microlattices. Strength y is
defined as the collapse stress of the lattice from uniaxial loading. Measurements are shown for (A) a solid HDDA
polymer lattice (y = 3.8 MPa) and (B) a Ni-P hollow-tube metallic lattice (y = 90 KPa).
Fig. S8
Stress-strain curves obtained from cyclic uniaxial compression of Al2O3 octet-truss microlattices. (A) Catastrophic
collapse was observed at first loading in the solid Al2O3 lattice with a density of 386 kg/m3. (B) and (C) For hollowtube Al2O3 lattices with mass densities of 10.2 kg/m3 and 1.7 kg/m3 , respectively, compression was performed to
progressively larger stress with complete unloading between each cycle. Discontinuities which are attributed to local
collapses in are indicated by arrows in the stress strain curves.
Fig. S9
A weight-bearing comparison between stretch- and bend-dominated lattices. 4x4x4 mm3 solid polymeric
microlattices, both fabricated from the same PEGDA material with a relative density of 5%, carrying at least
160,000 times their own weight. The octet-truss microlattice at left demonstrates markedly superior specific
stiffness compared with the open-cell foam (Kelvin foam) at right.
Table S1.
Summary of geometric and material properties for selected microlattices fabricated in this work.
Movie S1
Uniaxial incremental cyclic compression of a 5% solid polymer (HDDA) bend-dominated lattice
(played at 4x speed). Compression was actuated using a precision air-bearing three-axis
positioning system (Aerotek, Inc., Pittsburgh, PA). In the test configuration, the bottom surface
remained fixed while the top surface was displaced in the vertical direction. Each successive
compression cycle increases the lattice deformation by approximately 10% strain.
Movie S2
Uniaxial incremental cyclic compression of a 9.8% solid polymer (HDDA) stretch-dominated
lattice (played at 4x speed). The equipment and test configuration are the same as Movie S1.
Each successive compression cycle increases the lattice deformation by approximately 15%
Movie S3
Uniaxial cyclic compression of a stretch-dominated ultralight hollow-tube Ni-P lattice (=11.8
kg/m3) with increases in lattice strain by approximately 10 µm at each subsequent loading cycle.
The equipment and test configuration are the same as Movie S1.
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