three-dimensional printing of complex-shaped alumina/glass composites
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DOI: 10.1002/adem.200900213MM
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Three-Dimensional Printing of Complex-Shaped Alumina/Glass Composites**
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By Wei Zhang*, Reinhold Melcher, Nahum Travitzky, Rajendra Kumar Bordia and Peter GreilN
Alumina/glass composites were fabricated by three-dimensional printing (3DPTM) and pressurelessinfiltration of lanthanum-alumino-silicate glass into sintered porous alumina preforms. The preformswere printed using an alumina/dextrin powder blend as a precursor material. They were sintered at1600 8C for 2 h prior to glass infiltration at 1100 8C for 2 h. The influence of layer thickness and sampleorientation within the building chamber of the 3D-printer on microstructure, porosity, and mechanicalproperties of the preforms and final composites was investigated. The increase of the layer thicknessfrom 90 to 150mm resulted in an increase of the total porosity from �19 to �39 vol% and thus, in adecrease of the mechanical properties of the sintered preforms. Bending strength and elastic modulus ofsintered preforms were found to attain significantly higher values for samples orientated along theY-axis of the 3D-printer compared to those orientated along the X- or the Z-axis, respectively.Fabricated Al2O3/glass composites exhibit improved fracture toughness, bending strength, Young’smodulus, and Vickers hardness up to 3.6MPa m1/2, 175MPa, 228GPa, and 12GPa, respectively.Prototypes were fabricated on the basis of computer tomography data and computer aided design data toshow geometric capability of the process.
[*] W. Zhang, Dr. N. Travitzky, Prof. P. GreilDepartment of Materials Science, Glass and CeramicsFriedrich-Alexander-University Erlangen-NurembergMartensstr. 5, 91058 Erlangen, GermanyE-mail: [email protected]
R. MelcherRobert Bosch GmbH Bamberg Plant, Diesel SystemsPostfach 1160, 96045 Bamberg, Germany
Prof. R. K. BordiaDepartment of Materials Science and Engineering, Universityof WashingtonRoberts Hall, Box 352120, Seattle, WA 98195, USA
Dr. N. Travitzky, Prof. P. GreilCentre for advanced Materials and Processes Friedrich-Alexander-University Erlangen-NurembergDr. -Mack-Strasse 81, 90762 Fuerth, Germany
[**] The authors thank the Deutsche Forschungsgemeinschaft(DFG) for financial support and Dr. M. Stephan of VitaZahnfabrik, Bad Saeckingen, Germany for supply of glasspowder and kind advice. R. K. Bordia thanks to the Alexandervon Humboldt Foundation for financial support.
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Alumina ceramics are well studied and universally used
low cost materials possessing such attractive properties as
excellent wear and oxidation resistance, good high tempera-
ture strength, etc.[1] In particular, alumina/glass composites
are attractive materials for a wide range of engineering
applications or biomedical uses such as dental restorations
due to their high strength, low thermal conductivity, abrasion
resistance, biocompatibility, and esthetics.[2–6] Usually, alu-
mina/glass composites are fabricated using an approach
where alumina preforms are prepared by traditional pro-
cesses such as cold pressing or slip casting, followed by
continuous partial sintering at 1100–1400 8C and subsequent
glass melt infiltration into the porous structure at
950–1100 8C.[2,4,7,8] Nevertheless, there still remains the
challenge of forming complex-shaped alumina/glass compo-
sites as the above-mentioned process is limited in respect to
the required geometric variety of parts.[2,9] Solid-free
form (SFF) or rapid prototyping (RP) techniques, offer
completely new paths in order to realize complex-shaped
ceramic bodies.
Three-dimensional printing (3DPTM) is an RP process in
which powdered material is deposited in layers.[10] A binder
solution is locally applied on a powder layer by an ink-jet print
head, causing the powder particles to bind to one another and
to the printed cross-section one level below. This process is
repeated until the entire part is printed. A full dense and
complex-shaped composites were fabricated by 3DP process
and subsequent pressureless melt infiltration.[11–16]
The processing parameters of 3D-printing strongly influ-
ence the microstructure and mechanical properties of the
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Fig. 1. (a) Process diagram for fabrication of alumina/glass composites and (b) sche-matic of building chamber: machine directions and sample orientations.
printed bodies. For instance, a binder saturation determined
the geometry of an elementary building units, resulting from a
single ink drop in the powder bed and their interaction with
the powder bed, through which the surface finish and
microstructure of the preform may be controlled.[11] In addition,
the preforms printed with slower printing speed showed better
surface structure compared to those with faster speeds even
though the binder saturation levels were similar.[13,17] Influence
of the layer thickness and printing direction on microstructures
and mechanical properties of printed bodies, however, has not
been reported up to date in detail.
The purpose of the present work was to study the effect of
layer thickness and printing direction on the microstructure
and mechanical properties of alumina/glass composites
fabricated by 3D-prnting of porous Al2O3 preforms followed
by their sintering and pressureless melt infiltration with
lanthanum-aluminosilicate glass.
Experimental
Powder Processing
Granulate powder for the fabrication of porous ceramic
preforms was prepared as follows: 120 g dextrin powder
(Superior Gelb/mittel-F, Suedstaerke, Schrobenhausen, Ger-
many) as binder and 19 g of dispersant (Dolapix A88,
Zschimmer & Schwarz, Lahnstein/Rhein, Germany) were
dissolved in 1.35 L distilled water. Subsequently, 1786 g
a-Al2O3 powder (CT 3000 SG, Almatis, Ludwigshafen,
Germany) and 94 g g-Al2O3 powder (PG feinst, Almatis) with
average particle sizes of 0.8 and 3.4mm, respectively, were
suspended. The mixture was homogenized by tumbling it in a
2 L polyethylene bottle with alumina grinding balls for 48 h.
The slurry was freeze dried at 50 8C/37 Pa (Delta 2–24, Christ,
Osterode/Harz, Germany). The dry batch was ball-milled in a
jar for 72 h and sieved through 150mm mesh.
3D-Printing
3D-printing was carried out on a Z310 printer
(Z-Corporation, Burlington/MA, USA) using a water-based
printing solution (ZB56, Z-Corpation). Two types of samples
were fabricated. Bar-shaped samples with nominal dimen-
sions of 6� 7� 60 mm3 were printed for the purpose of
mechanical testing of sintered preforms prior to infiltration.
For infiltration with the glass, plate-shaped samples with a
nominal dimension of 5� 72� 80 mm3 were printed. The
printed bodies were oriented in X-, Y-, and Z-orientations, the
layer thickness was adjusted to 90, 100, 120, and 150mm,
respectively, for each orientation. The orientations within the
build piston are defined as follows: X-orientation is the
direction of print head travel, Y-orientation is the direction of
gantry travel, Z-orientation is the direction of piston move-
ment (Fig. 1(b)). The printed samples were dried in the
powder bed at room temperature for at least 24 h.
Sintering and Melt Infiltration
The dried bar and plate-shaped samples were then sintered
in a resistance-heated furnace (HT 08/17, Nabertherm,
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Lilienthal, Germany) in air according to the following
temperature setup based on dilatometric results: 25–600 8Cat 600 8C h�1, 600–1100 8C at 1200 8C h�1, 1100–1600 8C at
300 8C h�1, holding time 2 h, 1600–25 8C at 600 8C h�1.
For infiltration, the 50 g lanthanum-aluminosilicate glass
powder (InCeram Alumina, VITA Zahnfabrik, Bad Saeckin-
gen, Germany) was mixed with 80 mL distilled water, 20 mL
glycerin, and 2.4 g dextrin to form a thin slurry which was
applied to one side of the porous alumina preform. The
amount of glass required was estimated on the basis of open
porosity determined by Hg-porosimetry. The plates with
coated face upward were dried and then heated to 1100 8C in
air at 1000 8C h�1 resistance-heated furnace, holding time 2 h.
The cooling was accomplished as follows: 1100–650 8C at
450 8C h�1, 650–400 8C at 170 8C h�1, 400–25 8C at 300 8C h�1. A
process overview is given in Figure 1(a).
Characterization
Particle size distribution of the granulate powder was
measured in an air stream by laser granulometry (Mastersizer
2000, Malvern Instruments, Malvern, Great Britain). The
porosity and pore size distribution of sintered alumina
preforms was measured by Hg-porosimetry (Pascal 140,
Thermo Electron, Rodano/Milan, Italy). Scanning electron
microscopy (SEM, Quanta 200, FEI, Prague, Czech Republic)
was applied for microstructural and fractographical analysis.
The samples for SEM analysis were ground and polished to a
6mm diamond finish and sputtered with gold. The phase
analysis of the infiltrated alumina samples was conducted by
X-ray powder diffraction (XRD) using monochromatic Cu Ka
radiation at a scanning rate of 0.75 min�1 over a 2u range of
10–708 (D500, Siemens, Karlsruhe, Germany).
Bending strength and fracture toughness of sintered and
infiltrated samples with dimensions of 3� 4� 50 mm3 were
measured by four-point bending method using a universal
testing machine (Instron 4204, Instron Corporation, Canton/
MA, USA) with a crosshead speed of 0.5 mm min�1. The
average values of flexural strength were determined from
measurements conducted on at least 10 testing bars. The bars
were loaded with spans of 20 and 40 mm at room temperature.
The tensile surfaces of the samples were polished to a 6mm
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Fig. 3. Linear shrinkage in three dimensions of preforms after sintering in dependenceon layer thickness.
diamond finish prior to bending. Fracture toughness of
infiltrated samples was evaluated using the single-
edge-v-notched-beam (SEVNB) method. A saw cut was
tapered using a razor blade with 3mm diamond paste. The
overall depth of the notch, which was determined by light
microscopy (Leica M 420, Leica, Heerbrugg, Switzerland), was
�0.4mm. Young’s modulus of preforms and infiltrated samples
were calculated from the longitudinal sound propagation
velocity measurements using the impulse excitation technique
(Buzz-o-sonic, BuzzMac Software, Glendale, USA) [18]. The
Vickers hardness was conducted on the polished surfaces of
infiltrated specimens at a load of 49N with a duration of 15 s
using a diamond indenter (Zwick 3231, Zwick, Ulm, Germany).
The average hardness value was determined from 10 indentation
measurements.
Results and Discussion
Particle size distribution of (aþ g)-alumina/dextrin pow-
der is presented in Figure 2(a). This powder exhibits a
multimodal particle size distribution with particle size
ranging from 0.1 to 300mm. Representative pore size
distributions are shown for green and corresponding sintered
samples printed with layer thickness of 100mm in Figure 2(b).
The pore size distribution of printed green samples exhibits a
bimodal distribution, showing maxima at 0.3 and 45mm
diameter. The fine pores were eliminated or closed during
Fig. 2. (a) Particle size distribution of (aþ g)-alumina/dextrin powder blend and (b)pore size distribution in green and sintered preform (layer thickness 100mm).
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sintering. Therefore, only pores in the range of 40mm were
obtained in the preform after sintering.
Figure 3 shows linear shrinkages in three dimensions of
preforms after sintering in dependence on layer thickness.
No orientational or layer thickness dependences could be
observed considering the standard deviation. The linear
shrinkage of 18.7� 0.3% was measured for all orientations
and layer thicknesses.
An increase of volume fraction of porosity in the sintered
Al2O3 preform was obtained with the increase of the layer
thickness. A minimum porosity of 19� 0.5 vol% was mea-
sured at a layer thickness of 90mm and a maximum of
39� 2.1 vol% at a layer thickness of 150mm (Fig. 4).
Consequently, porosity, and thus layer thickness, has an
effect on the mechanical properties of the sintered preforms.
For example, bending strength increases from 29� 11 to
98� 6 MPa and Young’s modulus increases from 55� 24
to 178� 23 GPa as the layer thickness decreases from 150 to
90mm (Fig. 4). Increasing of the mechanical properties may
be explained by decreasing of porosity in layers with low
thicknesses (Fig. 4).
In addition, bending strength as well as Young’s modulus
of sintered samples is affected by their orientation within the
building chamber. In general, samples orientated in the Y-axis
exhibit significantly higher mechanical properties than
samples printed in the X- or Z-orientation (Fig. 4). This effect
is all the more pronounced for increasing layer thickness.
Average values measured for Z-orientation tend to exhibit
significantly higher deviations than those of X- and
Y-orientations, respectively. This behavior might be explained
by the formation of continuous strips along the Y-axis during
printing. In comparison, layers along the X-axis are composed
of those strips being joined to each other. Along the Z-axis,
single layers are laminated, exhibiting interfaces and alter-
nating porosity.
In the case of glass infiltrated samples, no distinct
correlation could be observed as differences between average
values of bending strength and Young’s modulus are smaller
and values of standard deviation tend to sample out higher:
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Fig. 4. (a) Effect of layer thickness on the porosity and bending strength of sintered andinfiltrated alumina preforms in dependence on sample orientation and (b) effect of layerthickness on Young’s modulus of sintered and infiltrated alumina preforms in depen-dence on sample orientation.
Fig. 6. Fracture toughness of sintered and infiltrated samples depending on layerthickness.
bending strength increases from 139� 17 to 175� 13 MPa as
the layer thickness decreases from 150 to 90mm; Young’s
modulus is �228� 15 GPa. The same effect is observed
considering hardness of sintered and infiltrated samples:
Vickers hardness of sintered preforms increases from 0.5� 0.2
to 3.2� 1.1 GPa as the layer thickness decreases from 150 to
90mm. In the case of infiltrated samples, Vickers hardness of
11.6� 3 GPa was measured (Fig. 5).
Fig. 5. Hardness HV-5 of sintered and infiltrated samples depending on layer thicknessand tested plane.
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Fracture toughness of sintered samples is also affected by
their orientation within the building chamber (Fig. 6). Sintered
samples orientated in the Y-axis exhibit significantly
higher fracture toughness than those printed in the X- or
Z-orientation; infiltrated samples orientated in the Y- and
X-axis exhibit significantly higher fracture toughness
(�3.6� 0.2 MPa m1/2) than samples printed in the
Z-orientation (�3� 0.2 MPa m1/2).
SEM examination of the alumina/glass composites
revealed that the glass phase wets the alumina preform well.
The glass melt infiltration into Al2O3 preforms resulted in a
nearly dense composite with homogeneous structure
(Fig. 7(a)). Some closed porosity remained within the struts
and cannot be filled with glass melt. Phase composition
examined by XRD reveals a-Al2O3 to be the only crystalline
phase.
The SEM micrograph of a fracture surface of the alumina/
glass composite shows that a large variety of fracture
mechanisms is present (Fig. 7(b)): a transgranular fracture
path running through the alumina grains (marked by
arrow 1); an intergranular fracture path surrounding the
grains along grain boundaries and along the glass/alumina
interface (marked by arrow 2); a fracture path through the
glass (marked by arrow 3), and crack bridging including
grain-pull-out. Due to residual, internal stress fields in the
Fig. 7. (a) SEM micrograph of alumina/glass composite (layer thickness 120mm) and(b) fracture surface: transgranular fracture path (arrow 1), intergranular fracturepath along the glass/alumina interface (arrow 2), and fracture path through the glass(arrow 3) in the alumina/glass composite.
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Fig. 8. Parts fabricated using the three-stage process: (a) sintered and infiltratedhalf-skull and (b) sintered and infiltrated turbine wheel.
glass phase, which interact with the stress field of the
propagating crack, the crack is partially forced to propagate
through the alumina grains instead of along the grain
boundaries or alumina/glass interfaces. Crack deflection on
the alumina grains is indicated by a non-planar crack along
grain boundaries and interfaces. Every change in direction
contributes to the crack deflection and increases the fracture
resistance.[19]
The geometric capability of the process can be displayed by
means of prototypes from the engineering biomedical fields
fabricated by the 3D-printing, sintering, and glass infiltration
process (Fig. 8). Parts with complex geometries on the base of
computer tomography (CT) data such as a half skull (Fig. 8(a))
and computer aided design (CAD) data such as a turbine
wheel (Fig. 8(b)) were fabricated. All the geometric details of
the data models were exactly reproduced in the final parts.
Conclusions
Nearly dense parts of alumina/lanthanum-glass composite
with complex geometries were successfully fabricated by
3D-printing, sintering of the printed bodies, and post-
pressureless glass infiltration into sintered porous preforms.
The relationship of layer thickness of the 3D-printing process,
porosity, and mechanical properties of alumina preforms to
those of the corresponding composites was clarified: in
general, the minimization of the layer thickness will reduce
the porosity, and, as a result, bending strength and Young’s
modulus of the alumina preforms increase. Sample orienta-
tion also has a significant influence: mechanical properties of
sintered preforms orientated along the Y-axis of 3D-printer
are higher than those along the X- and Z-axis, respectively.
The mechanical properties of alumina/glass composites were
considerably improved by infiltration with lanthanum-
ADVANCED ENGINEERING MATERIALS 2009, 11, No. 12 � 2009 WILEY-VCH Verl
aluminosilicate glass and tend to no longer depend on the
layer thickness and direction.
Received: July 16, 2009
Final Version: August 10, 2009
Published online: November 27, 2009
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