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Fabrication and in vitro characterization of three-
dimensional organic/inorganic scaffolds by robocasting
J. Russias, E. Saiz, S. Deville, K. Gryn, G. Liu, R.K. Nalla, A.P. Tomsia
Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720
Received 22 August 2006; revised 15 October 2006; accepted 3 November 2006
Published online 00 Month 2007 in Wiley InterScience (www.interscience.wiley.com). DOI: 10.1002/jbm.a.31237
Abstract: A key issue for the fabrication of scaffolds for tis-
sue engineering is the development of processing techniques
flexible enough to produce materials with a wide spectrum
of solubility (bioresorption rates) and mechanical properties
matching those of calcified tissues. These techniques must
also have the capability of generating adequate porosity to
further serve as a framework for cell penetration, new bone
formation, and subsequent remodeling. In this study we
show how hybrid organic/inorganic scaffolds with con-
trolled microstructures can be built using robotic assisted
deposition at room temperature. Polylactide or polycaprolac-
tone scaffolds with pore sizes ranging between 200–500 lm
and hydroxyapatite contents up to 70 wt % were fabricated.
Compressive tests revealed an anisotropic behavior of the
scaffolds, strongly dependant on their chemical composition.
The inclusion of an inorganic component increased their stiff-
ness but they were not brittle and could be easily machined
even for ceramic contents up to 70 wt %. The mechanical
properties of hybrid scaffolds did not degrade significantly
after 20 days in simulated body fluid. However, the stiffness
of pure polylactide scaffolds increased drastically due to
polymer densification. Scaffolds containing bioactive glasses
were also printed. After 20 days in simulated body fluid they
developed an apatite layer on their surface.
2007 Wiley
Periodicals, Inc. J Biomed Mater Res 79A: 000–000, 2007
Key words: freeform fabrication; scaffold; composite; deg-
radation; mechanical properties
INTRODUCTION
The demand for biomaterials to assist or replace
organ functions and improve quality of life is rapidly
increasing.1 Traditional biomaterials for bone replace-
ment are developed from materials designed origi-
nally for engineering applications that have serious
shortcomings associated to the fact that their physical
properties do not match those of the surrounding tis-
sue and, unlike natural bone, cannot self-repair or
adapt to changing physiological conditions. Thus, an
ideal solution, and a scientific research challenge, is to
develop bone-like biomaterials (or tissue engineering
scaffolds) that will be treated by the host as normal
tissue matrices and will integrate with bone tissue
while they are actively resorbed or remodeled in a
programmed way, with controlled osteogenic activity.
This material will require an interconnected pore net-
work with tailored surface chemistry for cell growth
and penetration, and the transport of nutrients and
metabolic waste. It should degrade at a controlled
rate matching the tissue repair rates producing only
metabolically acceptable substances and releasing
drugs or stimulating the growth of new bone tissue at
the fracture site by slowly releasing bone growth fac-
tors (e.g., bone morphogenic protein or transforming
growth factor-b) throughout its degradation process.
In addition, its mechanical properties should match
those of the host tissues and the strength and stability
of the material–tissue interface should be maintained
while the material is resorbed or remodeled.
In recent years there has been an increasing interest
in the fabrication of porous three-dimensional struc-
tures with complex functionalities not only for use in
tissue engineering but also for many other applica-
tions such energy generation, structural, electronics,
etc. Techniques such as solvent casting, particulate
leaching, gas foaming, etc.2–5 have been traditionally
used for synthesizing porous structures, but they suf-
fer from numerous drawbacks including need for
specialized tooling and molds, poor reproducibility,
lack of proper control of porosity and interconnectiv-
ity, and, consequently, poor, unpredictable mechani-
cal properties. This has been one of the motivations
behind the development of new solid freeform fabri-
cation techniques such as direct ink-jet printing,
robotic assisted deposition or robocasting, and hot-
melt printing, which usually involve ‘‘building’’ struc-
tures layer-by-layer by deposition of colloidal inks
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Correspondence to: E. Saiz; e-mail: esaiz@lbl.gov
Contract grant sponsor: National Institutes of Health;
contract grant number: 5R01 DE015633; Contract grant
sponsor: Department of Energy; contract grant number:
DE-AC03-76SF00098
' 2007 Wiley Periodicals, Inc.

following a computer design; this can be achieved
with great precision and reproducibility.6–11
One of the techniques that is gaining popularity
during the last few years is robocasting12,13 where
computer-controlled deposition of a thick slurry is
performed to form three-dimensional structures
layer-wise at room temperature. Until now, robocast-
ing has been used to print ceramic inks where consol-
idation of the structures is achieved through a fluid
through-gel transition during printing.14 One of the
great challenges is to develop the right inks or sus-
pensions to print materials with a wider range of
chemistries with precision and reproducibility. In this
work we demonstrate how this technique can be used
to build three-dimensional organic/inorganic hybrid
structures with controlled porosity, custom composi-
tions, and appealing mechanical properties. The inor-
ganic component of the materials is either hydroxy-
apatite (HA), an osteoconductive calcium phosphate
closely related to the inorganic component of bone, or
a bioactive silica-based glass. Since the formulation of
the first bioactive glasses by Hench, they have been
extensively used in the fabrication of biomaterials
and composites due to their capacity to generate hy-
droxyapatite and form excellent bond with osseous
tissue.15 For the organic component we have chosen
either polylactide (PLA) or polycaprolactone (PCL),
two biocompatible and fully resorbable polymers
with different stiffness, approved by the Food and
Drug Administration (FDA) for medical applications.
The work describes the development of hybrid inks
for the fabrication of porous scaffolds by robotic
assisted deposition and analyzes the physicochemical
factors that control their final properties, in particular
their in vitro evolution in simulated body fluid.
MATERIALS AND METHODS
Scaffolds fabrication
The inorganic component of the hybrid inks was either
commercially available HA powders (Trans-tech Adams-
town, MD, USA) with a particle size between 1 and 3 lm
[Fig. F11(a)], Bioglass
1 with the composition 45S5 developed
by Hench or a high-silica bioactive glass originally devel-
oped in our laboratory for the fabrication of coatings on
metallic alloys (6P53B)16 (Table T1I). The glass powders have
a wide particle size distribution with an average of *13 6
2 lm16 [Fig. 1(b)]. Properties of the different materials used
(density, Young’s modulus, and compressive strength) are
given in Table T2II.
Two different polymers were used in the inks: polylac-
tide (PLA) (molecular weight ¼ 92.1 kg/mol, 86.4% L iso-
mer) and polycaprolactone (PCL) (Sigma Aldrich, Saint-
Louis, MO, USA, molecular weight ¼ 80 kg/mol). To pre-
pare the inks, 5 g of polymer (PLA or PCL) were dissolved
in 15 mL of methylene chloride (J.T. Baker, Phillipsburg,
NJ, USA) at room temperature for 2 h using a magnetic
stirrer. The required amount of HA or glass was added to
the solution, together with 5 mL of denatured ethanol
(water content <0.1%) to control the viscosity of the slurry
and the evaporation kinetics and homogenized in a ball
mill with alumina balls for 1 h. The final quality of the ink
was assessed in terms of printability-measured as the min-
imum tip diameter suitable to extrude the ink without
clogging and stability (i.e., shape retention capacity during
drying and sintering) of the assembled structures.
Porous scaffolds were printed with a robocasting ma-
chine (3D inks, Stillwater, OK, USA) whose 3-axis motion
was independently controlled by a custom-designed, com-
puter aided direct-write program (Robocad 3.0, 3D inks,
Stillwater, OK, USA). The deposition process was carried
out in ambient atmosphere at room temperature. The three-
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Figure 1. Scanning electron micrograph of (a) hydroxyapatite powders and (b) 6P53B glass powders used in this work.
TABLE I
6P53B and Bioglass Compositions in Weight and mol %
SiO2 Na2O K2O CaO MgO P2O5
6P53B 52.7 (51.9) 10.3 (9.8) 2.8 (1.8) 18.0 (19.0) 10.2 (15.0) 6.0 (2.5)
Bioglass 45 (46.1) 24.5 (24.3) 0.0 (0.0) 24.5 (26.9) 0.0 (0.0) 6.0 (2.6)
Values in parentheses are in mol %.
2 RUSSIAS ET AL.
Journal of Biomedical Materials Research Part A DOI 10.1002/jbm.a

dimensional periodic scaffolds (*15 3 15 mm2 in length
and *4 mm in height corresponding to 17 layers) con-
sisted of a linear array of parallel rods in each layeraligned
such that their orientation was orthogonal to the previous
layer [Fig.F2 2(a)]. The center-to-center rod spacing was var-
ied between 0.5 and 1 mm in order to change the porosity
of the samples. Once a layer was printed, the nozzle was
raised by a fixed height, which depends on the tip diame-
ter, and another layer was deposited. The diameter of the
printing nozzles was varied from 5 to 410 lm and the
printing speed between 5 and 20 mm/s. During printing,
the flow rate was adjusted to the nozzle diameter and the
printing speed [Eq. (1)] in order to print a continuous line
of uniform thickness:
Vp¼
/t
/s

2
Vw
where Vp is the speed of the piston, /t is the diameter of
the tip, /s is the diameter of the syringe, and Vw is the
printing speed. Samples were printed on glass slides and
were easily removed after drying overnight at room tem-
perature in air.
Physical, microstructural, and mechanical
characterization
The average porosity was calculated using the measured
diameter of the printed rods and their spacing. The micro-
structure and composition of the printed parts were ana-
lyzed by X-ray diffraction (D500 diffractometer, Siemens
AG, Munich, Germany), optical microscopy (Axiotech
microscope, Carl Zeiss AG, Oberkochen, Germany), and
environmental-scanning electron microscopy (ESEM: S-
4300SE/N, Hitachi, USA) with associated energy disper-
sive spectroscopy (EDS). To obtain a three-dimensional
perspective of the structure, particularly to reveal subsur-
face defects, synchrotron X-ray computed tomography was
performed at the Advanced Light Source (ALS, Berkeley,
CA). Imaging was achieved on representative structures
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TABLE II
Density, Young’s Modulus, and Compressive Strength of the Raw Materials
Density (g/cm3) Young Modulus (GPa) Compressive Strength (MPa)
PLA 1.24 2.717 40–120 (pellet)18
PCL 1.14 0.417 –
Sintered dense HA 3.16 35–12019 120–90019
Bioglass1 2.7 35 *50018
6P53 B 2.716 7020 *7021
Figure 2. (a) Schematic representation of the robocasting process. The syringe displacement and the liquid flow through
the nozzle are controlled by the computer. (b) During the printing, the inks swell after leaving the capillaries and the
evaporation of the solvent creates a solid skin favoring the consolidation of the printed line. (c) Low magnification picture
showing one of the scaffolds fabricates in this study (HA/PLA, 70 wt % HA scaffold with 17 layers). (d) Three dimen-
sional reconstructed image of a HA/PLA (70 wt % HA) grid obtained by synchrotron X-ray computed tomography. It can
be observed that the material does not exhibit large defects. [Color figure can be viewed in the online issue, which is avail-
able at www.interscience.wiley.com.]
FABRICATION OF THREE-DIMENSIONAL ORGANIC/INORGANIC SCAFFOLDS AQ13
Journal of Biomedical Materials Research Part A DOI 10.1002/jbm.a

with 26 keV monochromatic X-rays and a 6 lm voxel size
(resolution). The tomography data were reconstructed into
three-dimensional images by a Fourier-filtered back-projec-
tion algorithm as described in detail elsewhere.22
The Vickers micro hardness of the printed lines was
measured by placing at least five indentations on 6 lm
polished surfaces with a load of 30 g. Compression tests of
strength, in the direction parallel and perpendicular to the
printing plane, were carried out on a ELF1 3200 series
voice-coil mechanical testing machine (EnduraTEC Inc.,
Minnetonka, MN) with a crosshead speed of 0.2 mm/min
on pieces (4 3 3 3 3 mm3) cut with a blade from the
printed scaffolds. At least four tests were performed for
each composition and each direction.
The in vitro response of the scaffolds was studied by
immersing the samples in 30 mL of simulated body fluid
(TableT3 III) at 378C for 20 days. The solution was prepared
by dissolving reagent-grade chemicals of NaCl, NaHCO3,
KCl, K2HPO4, MgCl2.6H2O, CaCl2, and (CH2OH)3CNH2
into distilled water and buffered with HCl to pH 7.25 at
378C.23 The evolution of the pH in the solution was meas-
ured with a pH meter (model 3000, VWR Scientific). After
20 days of immersion, the samples were washed with
large amounts of distilled water and air dried prior to
microstructural and mechanical characterization.
Thermal properties of PLA (in as received conditions,
after 20 days at 378C in air and after 20 days at 378C in
SBF) were determined using a differential scanning calo-
rimeter (DSC). The measurements were run under Neon/
Helium 50/50 volume mixture gas at a heating rate of
1008C/min using a Perkin Elmer Diamond DSC calibrated
with indium. The typical sample weight was 1 mg. There
was no thermal treatment of the samples before the first
heating scan in order to preserve their history. The
reported glass transition temperatures (Tg) were taken as
the mid-point of the step transition based on first heating
scans from 25 to 1508C.
RESULTS AND DISCUSSION
Physical and microstructural characterization
of the printed samples
The process used in the preparation of the inks is
based on the dissolution of the polymers in methyl-
ene chloride, like the standard procedure used for the
preparation of biodegradable polymer and HA/poly-
mer microspheres for drug delivery and composite
fabrication.24,25 The two key differences are that no
polyvinyl alcohol or any other difficult to eliminate
and potentially toxic surfactants are used and that
small quantities of a second organic phase (ethanol)
are added in order to tailor the ink viscosity and dry-
ing rates to the printing conditions. During printing,
solvent evaporation creates a solid skin on the
extruded line immediately after it exits the tip [Fig.
2(b)]. This skin confers some degree of rigidity to the
printed line and allows the fabrication of stable three-
dimensional structures [Figs. 2(c,d)]. Consolidation of
the lines is achieved by matching the printing speeds
to the drying kinetics of the ink. For very slow print-
ing speeds (typically less than 5 mm/s), the paste will
dry up in the printing nozzle, thus clogging it; if the
printing speed is too fast (faster than 20 mm/s), the
line diameter is not homogeneous and the lines can
be discontinuous.
The printing behavior of the inks also depends on
their inorganic content. High particle loads can result
in large viscosities and poor printability. However, by
adjusting the amount of ethanol (the amount of etha-
nol vs. inorganics, i.e., ceramic powders, is *40 wt %)
it is possible to prepare hybrid scaffolds with inor-
ganic contents as large as 70 wt %, very close to the
mineral content of cortical bone. This is very impor-
tant since the ceramic content influences the stiffness,
a key mechanical parameter that should be matched
with the host tissue. Typical required Young’s modu-
lus (E) varies between 0.4 and 350 MPa for soft tissue
and cartilage and up to 10–1500 MPa for hard tissue.6
Particular attention should be paid to the control of
the evaporation of the organic component during ink
preparation. Because of the use of methylene chloride
as a solvent, drying is quite rapid and the ink proper-
ties need to be maintained in the range adequate for
printing by hermetically sealing the containers. How-
ever, because the main mechanism of consolidation is
drying, control of the rheological properties does not
need to be as strict as for ceramic inks.13 Furthermore,
during robocasting of ceramic parts different strat-
egies (such as printing in oil baths) are used to avoid
drying and the associated dimensional changes and
stresses during printing.26 In the process described in
this work, the viscoplastic behavior of the parts
allows them to sustain the stresses generated during
simultaneous printing and drying. This viscoplastic
nature permits further stretching of the printed line to
form thin polymer and polymer/ceramic fibers and
ribbons (down to 1 lm in diameter). Figure F33 shows
very thin polymer and hybrid threads fabricated in
this manner. It can be clearly observed that the mini-
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TABLE III
Ion Concentrations of the SBF used in this Work and of Human Plasma
Ion Concentration (mM)
Naþ Kþ Ca2þ Mg2þ Cl
HCO
3 HPO
2

4 SO
2

4
SBF 142.0 5.0 2.5 1.5 147.8 4.2 1.0 0.5
Human plasma 142.0 5.0 2.5 1.5 103.0 27.0 1.0 0.5
4 RUSSIAS ET AL.
Journal of Biomedical Materials Research Part A DOI 10.1002/jbm.a

mum achievable diameter of the thread is controlled
by the size of ceramic particles. Biodegradable fibers
can be used in drug delivery applications or in the
fabrication of meshes for implants and scaffolds.27
It was not possible to successfully prepare an ink
containing the Bioglass composition developed by
Hench. This is attributed to the highly hygroscopic
nature of this glass (related to its high bioactivity)
that promotes fast reactions with small amounts of
water in the environment or in the ink. This fact also
hampers the use of Bioglass in the preparation of sus-
pensions to coat metallic alloys by enameling28 and
motivated the development of a family of high-silica
bioactive glasses like 6P53B that are mush less hygro-
scopic and more amenable to the preparation of sus-
pensions. Although several techniques have been pro-
posed to prepare Bioglass/polymer foams and porous
materials18,29,30 by using the 6P53B composition, it
was possible to prepare hybrid inks optimized for
rapid prototyping.
Partial wetting of the ink on the printed material
increases the contact area and provides better bond-
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Figure 3. (a) Pure PLA flexible ribbon and (b) thin HA/PLA (30/70 wt %) thread extruded after printing. Hydroxyapa-
tite grains are embedded in the polymer fiber and determine its minimum diameter.
Figure 4. Scanning electron micrographs showing different aspects of the microstructure. All the scaffolds showed in this
picture have been printed using a nozzle with an internal diameter of 410 lm. (a) Three dimensional reconstructed image
of the junction between HA/PLA (70 wt % HA) printed lines obtained by synchrotron X-ray computed tomography. It
can be observed that the ink wets partially the printed lines providing additional support and leading to good bonding
between lines (white circles). (b) Close up of an HA/PLA (70 wt % HA) printed line showing the homogeneous distribu-
tion of the HA particles (white) in the polymer. (c) HA/PCL (70 wt % HA) scaffold (center-to-center rod spacing *0.5 mm,
corresponding to a porosity of *55%); the pore size and thickness of the printed lines is very homogeneous. (d) 6P53B
glass/PLA (70 wt % 6P53B glass) scaffold (center-to-center rod spacing *1 mm, corresponding to a porosity of *75%).
[Color figure can be viewed in the online issue, which is available at www.interscience.wiley.com.]
FABRICATION OF THREE-DIMENSIONAL ORGANIC/INORGANIC SCAFFOLDS AQ15
Journal of Biomedical Materials Research Part A DOI 10.1002/jbm.a

ing between lines [Fig.F4 4(a)], enhancing the stability of
the structure. In the hybrid scaffolds the ceramic
phase is homogeneously distributed in the polymer
[Fig. 4(b)] and there are no visible surface defects or
microporosity. By changing the center-to-center rod
spacing (0.5 and 1 mm), samples with 2 different
porosities are processed. A center-to-center rod spac-
ing of 0.5 mm corresponds to a porosity of 55% and a
pore size of about 200 lm [Fig. 4(c)], while a center-
to-center rod spacing of 1 mm gives a porosity of 75%
and a pore size of about 500 lm [Fig. 4(d)].
The final line thickness depends on several factors:
nozzle diameter, printing speed, printing height [DZ
in Fig. 2(b)], and drying shrinkage (typically *25% in
volume). After leaving the nozzle, the ink was
observed to swell, with the degree of swelling depend-
ing on the printing speed (flow rate) and the nozzle di-
ameter. As shown in Figure F55, for a 410-lm tip, the
swelling and the printing speed influence are minimal
and the drying shrinkage is important since the final
average line diameter is around 200 lm. For a very
thin capillary of 5 lm, there is a significant swelling
that can be controlled by increasing the printing speed
and the diameter of the printed lines can be reduced
from 400 to 100 lm, without affecting the pore size, by
changing the printing speed from 5 to 20 mm/s.
In a further step, the scaffolds can be infiltrated to
make dense samples with controlled composition and
phase distributions. By using phases with different
degradation rates porosity can be created ‘‘in situ’’
providing interesting possibilities for the control of
the bioresoption and the mechanical behavior in vivo.
For example, we have been able to successfully infil-
trate a PLA/HA (70 wt % of HA) scaffold with a
PCL/HA slurry (70 wt % of HA, with a slurry pre-
pared exactly in the same conditions than the inks for
printing) under vacuum (150 mbar) without leaving
residual porosity (Fig. F66). The systematic analysis of
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Figure 5. Evolution of the final line diameter versus the
printing speed for the printing of a pure PLA ink through
nozzles with two internal diameters. The insert shows a
scanning electron micrograph of a PLA grid printed with a
5-lm tip at a printing speed of 15 mm/s.
Figure 6. Scanning electron micrograph of a PLA/HA (70 wt % HA) scaffold cross section infiltrated with PCL/HA
(70 wt % HA). No residual porosity can be observed after the infiltration process.
6 RUSSIAS ET AL.
Journal of Biomedical Materials Research Part A DOI 10.1002/jbm.a

the microstructure and properties of these materials
will be the subject of a future work.
In vitro behavior in simulated body fluid
In vitro tests in cell-free solutions with ionic concen-
trations similar to those of body fluids allows analysis
of the chemical and microstructural evolution of the
materials under conditions that simulate their biologi-
cal interactions with the body and provide fundamen-
tal data to predict and understand their in vivo behav-
ior and long term stability.23 After 20 days in simu-
lated body fluid microporosity develops on the
surface of pure PLA scaffolds (pore size from 0.5 to
1 lm) (Fig.F7 7). To analyze the possible changes in the
polymer structure induced by the in vitro treatment at
378C, DSC analysis of the ‘‘as received’’ PLA and after
20 days at 378C in air and in SBF are performed (Fig.
F8 8). The DSC analysis confirms that, as expected, the
as received PLA is amorphous since it contains 14%
of D isomer. However, its glass transition tempera-
ture increases from 33.08C for the as received materi-
als to 49.18C after 20 days at 378C or to 60.88C after 20
days at 378C in SBF. These results indicate that the
structure of the polymer is changing during the in
vitro tests: the polymer chains are organizing and the
polymer is getting denser. This temperature and hu-
midity are two factors responsible for the densifica-
tion of the PLA. No diffusion of SBF ions into the
polymer is observed by EDS analyses. The volume
change associated with the densification might be re-
sponsible for the formation of the micropores. The
addition of HA to the PLA reduces significantly the
shrinkage of the material and micron size pores do
not form on HA/PLA scaffolds even after 20 days in
SBF. Degradation of the PCL-based scaffolds (with or
without hydroxyapatite) results in the formation of a
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Figure 7. (a) Scanning electron micrographs of a PLA scaffold after 20 days in SBF. Note the presence of micron size
pores homogeneously dispersed on the polymer surface. (b) The inset shows these pores at a larger magnification.
Figure 8. Evolution of the glass transition temperature
and stiffness up to the elastic limit for PLA scaffolds in an
as received condition, after 20 days at 378C in air and after
20 days at 378C in simulated body fluid.
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network of polymer fibers between the printed lines
(Fig.F9 9) while their surface becomes rougher.
The formation of acidic products during the hydro-
lytic degradation of resorbable polymer materials has
been reported to cause adverse body reactions and
additions of various calcium phosphates to the poly-
mer have been proposed as a mean to buffer the
release of acidic products and avoid these reac-
tions.31–34 FigureF10 10 shows the pH evolution of the
simulated body fluid solution in which the PLA-
based samples were immersed. The behavior is simi-
lar for PCL based scaffolds. The pH remains stable,
around 7.25, for pure polymer and HA/polymer com-
posites. The presence of glass particles has two
effects: the pH of the solution increases to reach 7.8
after 20 days of immersion for scaffolds containing
6P53B glass and apatite crystals precipitate on the
scaffold surface (Fig.F11 11). This is due to a rapid ion
exchange of Naþ from the glass with Hþ and H3O
þ
followed by a polycondensation reaction of surface
silanols to create a high-surface area silica gel. This
gel can provide a large number of sites for heteroge-
neous nucleation and crystallization of a biologi-
cally reactive hydroxy-carbonate apatite (HCA) layer
equivalent to the inorganic mineral phase of bone.15
The ion exchange from the bioactive glass can buffer
the acidic products resulting from polymer degrada-
tion and promote apatite precipitation. It has been
proposed that this growing apatite layer favors the
bonding to bone of bioactive glasses and has a signifi-
cant impact on the activity of osteogenic cells.35
Mechanical characterization
Microhardness
As expected from the intrinsic properties of the
polymers (Table II) the printed lines of the PLA-based
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Figure 9. Scanning electron micrograph of a PCL/HA (70 wt % HA) scaffold after 20 days in SBF. Polymer filaments are
bridging the lines. The degradation of the scaffold surface is clearly visible.
Figure 10. pH evolution of the solution during 20 days of
immersion in SBF for PLA, PLA/HA, and PLA/6P53B
glass materials. Note the increase in pH for the 6P53B
glass/PLA scaffold.
8 RUSSIAS ET AL.
Journal of Biomedical Materials Research Part A DOI 10.1002/jbm.a

scaffolds are much harder than the equivalent PCL
composites (Fig.F12 12). The hardness of the glass is
6.2 GPa.16 There is a wide spread in the hardness data
reported for sintered HA but it seems to range
between 3.5 and 6.5 GPa for a fully dense ceramic.36–38
The addition of glass increases the hardness but there
is a larger dispersion in the measurements probably
due to the larger grain size and the larger dispersion
of the grain size distribution. After immersion in
simulated body fluid the hardness of pure PLA and
PLA/HA materials increase slightly. This might be
due to the observed densification of the polylactide
during the in vitro tests.
Compression strengths
Compression tests are performed in two directions
(Fig.F13 13), perpendicular (direction 1) and parallel
(direction 2) to the printing plane. During the tests,
the scaffolds do not fail in a brittle manner and show
an elasto-plastic response with large plastic ‘‘yielding".
The insets on Figure 13 show SEM pictures of the
samples in both directions after compression. Only
PLA/HA samples tested in the direction 2 show a
maximum in the stress vs. strain curve at around 7–
10 MPa. This is probably due to the fact that they are
harder and stiffer and the printed lines buckle instead
of being continuously deformed like in the other scaf-
folds (Fig. 13).
The lines do not debond during testing in any
direction, indicating excellent adhesion between the
printed rods. The stress–strain curves for samples
with the same composition and microsctructure are
very similar (Fig. 13), allowing the comparison of the
mechanical response of different materials. Consider-
ing the samples’ geometry and due to their large de-
formation it is difficult to calculate an absolute value
of the Young’s modulus from the compression tests.
However, the slope of the compression curves (calcu-
lated using a linear fit of the data up to the elastic
limit) can provide several trends.
1. The porosity (in the range used in this work,
between 55 and 75 vol %) does not affect signifi-
cantly the mechanical response of the samples.
2. The mechanical response of the scaffolds is
clearly anisotropic (Fig. 13).
3. The mechanical properties can be easily
adjusted by controlling the composition of the
material: type of polymer and inorganic con-
tent. Addition of up to 70 wt % of hydroxyapa-
tite increases the stiffness up to two orders of
magnitude [Fig. F1414(a)] and PLA-based scaffolds
are much stiffer than those containing PCL
[Fig. 14(b)]. For an HA/polymer scaffolds with
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Figure 11. Scanning electron micrographs of the 6P53B glass/PLA scaffold after 20 days of immersion in SBF. (a) Apatite
crystals have precipitated on the scaffold surface. (b) The apatite layers consist of small nano-crystals (inferior to 50 nm)
that cover progressively the scaffold.
Figure 12. Micro hardness of the printed lines samples in
as received conditions and after 20 days in a SBF. Or-
ganic/inorganic samples have 70 wt % of inorganic phase.
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Journal of Biomedical Materials Research Part A DOI 10.1002/jbm.a

a porosity of 55% (pore size *200 3 200 lm2)
and an HA content of 70 wt %, the values of
the slopes in the direction perpendicular to the
printing plane varies between 84 6 9 MPa for
a PLA-based materials and 24 6 5 MPa for
PCL based materials. In the parallel direction,
the slopes vary between 150 6 40 and 110 6
20 MPa for PLA and PCL-based scaffolds,
respectively. This is expected since the elastic
modulus of dense PLA and PCL are respec-
tively 2.7 and 0.4 GPa (Table II).
4. After twenty days in SBF, pure PLA scaffolds
become much stiffer and the slope of the strain–
stress curve up to the elastic limit increases from
3 6 2 to 76 6 19 MPa [Fig. F1515(a)]. Like the paral-
lel increase in microhardness, this is related to
the observed densification of the polymer as a
result of the in vitro treatment. The PLA scaf-
folds also become stiffer after 20 days in air at
378C. The slope’s average value of their strain–
stress curve up to the elastic limit is 35 6 6 MPa.
As it is confirmed by the DSC analysis, the
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Figure 13. Strain–stress curves obtained from compression tests made on different grids of HA/PLA (70 wt % HA) sam-
ples with 55% of porosity in the two directions (a) perpendicular (direction 1) and (b) parallel (direction 2) to the printing
plane. The insets show some SEM images of the samples after compression tests in direction 1 and 2. The stress–strain
curves in each direction are very similar. The samples do not fail in a brittle manner and the behavior is anisotropic. Buck-
ling of the printed lines of the PLA/HA scaffolds can be clearly observed in the inset [Fig. 13(b)].
Figure 14. (a) Influence of the ceramic content on the compressive behavior of the scaffolds. The strain–stress curves are
obtained from compression tests made on PLA and HA/PLA samples in two directions. Addition of hydroxyapatite
increases significantly the stiffness of the material. (b) Influence of the organic phase. Samples containing PLA are stiffer.
10 RUSSIAS ET AL.
Journal of Biomedical Materials Research Part A DOI 10.1002/jbm.a

long treatment at 378C promotes densification
and this process is enhanced in SBF (Fig. 8).
The moisture plasticize the polymer chains
favoring their rearrangement and enhancing
densification. This result suggests that thermal
treatments can be used to further manipulate
the mechanical response of the material.
5. The compressive behavior of the PLA or PCL-
based hybrid scaffolds containing 70 wt % of
hydroxyapatite do not change significantly af-
ter 20 days in simulated body fluid. The elastic
modulus of HA is 10–100 times larger than the
one of the polymer (Table II) and the main con-
tribution to the Young’s modulus of the com-
posites comes from the inorganic phase that
does not undergo any visible degradation dur-
ing in vitro testing in SBF. However, the bioac-
tive glasses react with simulated body fluid
and there is an appreciable decrease of the
stiffness in the scaffolds containing 6P53B after
20 days in SBF [Fig. 15(b)].
SUMMARY
This work demonstrates how robotic assisted depo-
sition can be successfully used for the fabrication of
porous hybrid organic/inorganic materials of various
chemical compositions with well controlled architec-
ture and porosity. The technique is versatile enough
to allow the combination of a wide range of materials,
including bioactive glasses whose addition can be
used to buffer the possible formation of acidic degra-
dation product coming from the hydrolytic degrada-
tion of the polymer and to promote apatite formation.
In addition, by selecting the adequate organic compo-
nent and manipulating the organic/inorganic ratio of
the scaffolds, we can control their stiffness and fabri-
cate materials much more rigid than porous polymers
while avoiding the brittleness of ceramic parts. More-
over, multi-component materials can be fabricated
through the simultaneous use of multiple nozzles and
because the processing is performed at room temper-
ature, in situ seeding with cells and addition of drugs
or growth factors to the organic component is easily
achievable. These characteristics suggest that robotic
assisted deposition can be a fast and economical alter-
native for the fabrication of ‘‘on demand’’ scaffolds for
biomedical applications.
We acknowledge the support of the dedicated tomogra-
phy beamline (BL 8.3.2) at the Advanced Light Source
(ALS).
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