High strength mullite-bond SiC porous ceramics fabricated by digital light processing

Abstract

Fabricating SiC ceramics via the digital light processing (DLP) technology is of great challenge due to strong light absorption and high refractive index of deep-colored SiC powders, which highly differ from those of resin, and thus significantly affect the curing performance of the photosensitive SiC slurry. In this paper, a thin silicon oxide (SiO2) layer was in-situ formed on the surface of SiC powders by pre-oxidation treatment. This method was proven to effectively improve the curing ability of SiC slurry. The SiC photosensitive slurry was fabricated with solid content of 55 vol% and viscosity of 7.77 Pa·s (shear rate of 30 s⁻¹). The curing thickness was 50 μm with exposure time of only 5 s. Then, a well-designed sintering additive was added to completely convert low-strength SiO2 into mullite reinforcement during sintering. Complex-shaped mullite-bond SiC ceramics were successfully fabricated. The flexural strength of SiC ceramics sintered at 1550 ℃ in air reached 97.6 MPa with porosity of 39.2 vol%, as high as those prepared by spark plasma sintering (SPS) techniques.

Full text


High strength mullite-bond SiC porous ceramics fabricated by
digital light processing
JianSun1, JingdeZhang1,2, XuZhang1, ZiheLi3, JianzhangLi4, SijieWei1,5, WeibinZhang1, WeiliWang1,
GuifangHan1,
✉
1 Key Laboratory for Liquid-Solid Structural Evolution and Processing of Materials (Ministry of Education), School of Materials Science
and Engineering, Shandong University, Jinan 250061, China
2 Key Laboratory of Special Functional Aggregated Materials, Ministry of Education, Shandong University, Jinan 250100, China
3 Materials and Structures Research Centre, Department of Mechanical Engineering, University of Bath, Bath BA2 7AY, UK
4 National Engineering Research Centre of Ceramic Matrix Composite Manufacture Technology, Xi’an Golden Mountain Ceramic
Composites Co., Ltd., Xi’an 710118, China
5 School of Physics and Materials Science, Changji University, Changji 831100, China
Received: September 20, 2023; Revised: November 11, 2023; Accepted: November 27, 2023
©The Author(s) 2024. This is an open access article under the terms of the Creative Commons Attribution 4.0 International License
(CC BY 4.0, http://creativecommons.org/licenses/by/4.0/).
Abstract: FabricatingSiCceramicsviathedigitallightprocessing(DLP)technologyisofgreatchallengeduetostronglight
absorption and high refractive index of deep-colored SiC powders, which highly differ from those of resin, and thus
significantlyaffectthecuringperformanceofthephotosensitiveSiCslurry.Inthispaper,athinsiliconoxide(SiO2)layerwas
in-situformedonthesurfaceofSiCpowdersbypre-oxidationtreatment.Thismethodwasproventoeffectivelyimprovethe
curing ability of SiC slurry. The SiC photosensitive slurry was fabricated with solid content of 55 vol% and viscosity of
7.77Pa·s (shear rate of 30 s−1). The curing thickness was 50μm with exposure time of only 5 s. Then, a well-designed
sinteringadditive wasaddedto completelyconvertlow-strength SiO2intomullite reinforcementduringsintering. Complex-
shapedmullite-bondSiCceramicsweresuccessfullyfabricated.TheflexuralstrengthofSiCceramicssinteredat1550℃in
airreached97.6MPawithporosityof39.2vol%,ashighasthosepreparedbysparkplasmasintering(SPS)techniques.
Keywords: digitallightprocessing(DLP);SiCceramics;pre-oxidation;mullite-bondSiC;mechanicalproperties

1Introduction
Due to the advantages of high strength/hardness, high thermal
conductivity, low thermal expansion coefficient, and excellent
temperature/corrosion/wear resistance [1−5], silicon carbide (SiC)
ceramics are widely used in many fields including thermal
insulating, filter, catalyst carriers, biology, and aerospace [6−10].
These excellent properties of SiC ceramics are due to a strong
Si–C covalent bond. However, this results in difficulty in sintering.
In addition, the high strength and hardness of the SiC ceramics
make it difficult to machine or process [11,12], which limits the
complexity of SiC components and therefore their applications.
As a novel concept, the additive manufacturing technology can
realize direct fabrication of complex structure ceramics without
machining [13]. Among a range of additive manufacturing
techniques, since the digital light processing (DLP) technique has
high precision and relatively fast production speed, DLP has been
widely applied in the fabrication of complex ceramic components
[14−17]. The mechanism of photocuring is based on the
polymerization of the photosensitive resin under the irradiation of
specific wavelengths of ultraviolet (UV) light [18,19]. In the first
step, ceramic powders are added to this photosensitive resin to
form uniformly dispersed ceramic slurry. Then, the slurry is cured
layer by layer with UV light exposure, and simultaneously ceramic
particles are wrapped inside the cured layer to form the designed
shape. The DLP technique has been successfully applied in the
printing of Al2O3 [20−22], ZrO2 [23], and other oxide ceramics.
However, deep-colored SiC has strong light absorption and a
refractive index much higher than that of resin. This seriously
reduces the penetration depth of light, and the amount of light
that can be utilized by photosensitive resin, thereby decreasing the
curing thickness of SiC slurry [24].
To solve this problem, researchers decreased the solid content
of SiC inside the slurry and/or increased the particle size of SiC
powders [24,25]. However, the sample fabricated with low-solid-
content slurry showed large shrinkage after the remove of
polymer agents, which induced defects and thus poor mechanical
properties of the sintered samples. In addition, the large particle
size of SiC also resulted in low strength of fabricated parts.
Another method to increase the curing thickness of the SiC slurry
is to coat shell materials with low absorbance on the surface of SiC
particles by physical or chemical methods. As a result, the shell
layer can reduce the absorbance and improve curing ability
[26−29]. Interestingly, high-temperature oxidation treatment of
the SiC powders in air can generate in-situ SiO2 shell layers with
low absorbance, which has been proven to increase the curing
depth of the SiC slurry [26,27].
However, SiO2 has low mechanical properties and undergoes
phase change during the cooling process accompanied by volume
changes. The mismatch of the thermal expansion coefficient
(CTE) between cristobalite SiO2 and SiC (CTEcristobalite =

✉
Corresponding author.
E-mail: gfhan@sdu.edu.cn

Journal of Advanced Ceramics
2024, 13(1): 53−62
https://doi.org/10.26599/JAC.2024.9220835 ResearchArticle

https://doi.org/10.26599/JAC.2024.9220835


0.5×10−6 ℃−1, CTESiC = 4.9×10−6 ℃−1) [30] results in weak bonding
between them and therefore the poor performance of the SiC
ceramics. All these factors limit mechanical strength of the printed
SiC components. In our previous work [31], we sintered SiC
porous ceramics at 1550 ℃ in air with the strength comparable to
that fabricated from the spark plasma sintering (SPS) technique.
The oxidation-derived SiO2 from SiC at high temperatures in air
was converted into high-strength mullite bonding phases and
whisker reinforcement due to the rational design of the composite
sintering additives. In the Al(OH)3–Y2O3–CaF2 composite
additive, Al(OH)3 provides an Al2O3 source to consume SiO2, and
Y2O3 promotes the liquid mass transfer and the formation
reaction of mullite, while CaF2 introduces gaseous mass transfer
and enhances the formation of mullite whiskers. By converting
low-strength SiO2 to the mullite bond phase and introducing
mullite whisker reinforcement, this technique has been proven to
effectively enhance mechanical properties of the SiC ceramics.
Therefore, in this study, fine SiC powders with an average
particle size of 5.28 μm were pre-oxidized to form a thin SiO2
shell. The curing depth of the thus fabricated photosensitive slurry
was significantly increased, and complex-shaped SiC components
were successfully fabricated. SiO2 was completely converted to
mullite by carefully adjusting the added amount of
Al(OH)3–Y2O3–CaF2 sintering additives. The fabricated mullite-
bond SiC ceramics achieved high flexural strength comparable to
that fabricated by the SPS method. This work therefore provides a
feasible strategy to enhance the curing ability of the SiC slurry and
the mechanical property of sintered SiC at the same time, such
that we hope to excite the whole society to realize a low-cost and
precise photo-polymerization fabrication of complex-shaped deep-
colored ceramic materials such as SiC and Si3N4.
2Experimental
2.1Materials
Commercial SiC ceramic powders (purity 99.9%, d50 = 5.288 μm,
Nangong Ruiteng Alloy Material Co., Ltd., China) were chosen as
raw materials. Aluminum hydroxide (Al(OH)3, purity 99.9%,
average particle size of 10 μm, Aladdin Chemical Co., Ltd., China),
yttria (Y2O3, purity 99.9%, average particle size of 1 μm, Aladdin
Chemical Co., Ltd., China), and calcium fluoride (CaF2, purity
99.9%, average particle size of 1 μm, Shanghai Maclin Biochemical
Technology Co., Ltd., China) were used as sintering additives.
Tetramethyl ammonium hydroxide aqueous solution (TMAH,
25 wt%, AR, Aladdin Chemistry Co., Ltd., China) was used as the
dispersant (1.6 wt% of ceramic powders). Photosensitive resin
used in this paper was commercial resin containing monomers
and photoinitiators. The photoinitiator was 2,4,6-trimethylbenzoyl-
diphenyl phosphorus oxide (TPO). The monomer was a
dimethacrylate based substance, and the dispersant was KOS110
(Guangzhou Tai Runding New Materials Co., Ltd., China).
2.2Experimentalprocedure
The SiC powders were evenly spread into an Al2O3 crucible and
put into a chamber furnace for pre-oxidation in air atmosphere.
The heating rate was 10 ℃/min and dwelled at 1200 ℃ for 2, 4,
and 6 h. The pre-oxidized powders were ball milled at 300 r/min
for 2 h and then sieved through a 100-mesh sieve.
In our previous work [31], the amount of Al(OH)3 was
optimized as 6.4 wt% of that of the SiC powders. The amount of
Y2O3 was optimized as 1.6 wt% of that of SiC and Al(OH)3
powders. The amount of CaF2 was optimized as 2.0 wt% of that of
the Al(OH)3 powders. Since we implemented pre-oxidation in this
study and the debinding process in air also introduced oxidation
of SiC, the amount of Al(OH)3 was increased to 10, 12.5, and 15 wt%
to completely consume additional SiO2. Also, the amount of Y2O3
and CaF2 was also increased accordingly. All these powders were
put into anhydrous ethanol (analytically pure) to form a
suspension with solid loading of 50 vol%, and TMAH was added
to the suspension as the dispersant to enhance the mixing
uniformity of SiC and additive powders. After ball milling for 4 h
with a planetary ball mill at a rotating speed of 300 r/min, the
slurry was dried in a vacuum drying oven. In the process of
obtaining composite powders by drying, the drying temperature
was 90 ℃, and the time was more than 24 h. The dried powders
were then crushed and ground through a 100-mesh screen to
obtain composite powders.
The composite powders were mixed with commercial
photosensitive resin to obtain ceramic slurry with solid loadings of
45, 50, and 55 vol%, and KOS110 (5 wt% of the total weight of
ceramic powders) was used as the dispersant agent. After the
ceramic slurry was planetary ball milled for 5 h at a rotating speed
of 300 r/min, a vacuum drying oven (ZK-025, Shanghai
Experimental Instrument Factory, China) was used for deforming
of 1 h to obtain uniformly dispersed photocurable ceramic slurry.
The temperature of the ceramic slurry in the vacuum defoaming
process was controlled at room temperature of about 25 ℃.
The DLP printer (PC5003A-50, Xi'an Dianyun Biotechnology
Co., Ltd., China) with a UV light wavelength of 405 nm was
used in this study. The intensity of the light source used was
30 mW/cm2. This instrument used a digital mask device (DMD)
with a resolution of 1920 × 1080 pixels. The dot resolution on the
printing plane was 50 μm, and the control accuracy in the printing
direction was around 10 μm. For the printing parameters, the
single layer thickness and the UV light exposure time were set to
50 μm and 3 s, respectively. Printing was conducted layer-by-layer
until the designed shape was obtained. Furthermore, within the
scope of this investigation, ceramic green bodies were prepared
through the printing process and subsequently subjected to
sintering to yield the final ceramic products. The heating rate in
the debinding process was 1 ℃/min till 600 ℃. After holding for
1 h, the temperature was raised to 1200 ℃ at a rate of 5 ℃/min,
and finally to the sintering temperature of 1450, 1500, and 1550 ℃
at a heating rate of 10 ℃/min and holding time of 2 h.
2.3Characterizations
An X-ray diffractometer (D/max 2500 PC, Rigaku, Japan, Cu Kα,
λ = 0.1548 nm) was used to obtain X-ray diffraction (XRD)
patterns for sintered ceramic materials. A scanning electron
microscope (SEM; JSM-7800, JEOL, Japan) equipped with an
energy dispersion spectrometer (EDS) was used to study the
microstructure of powders and cross-sections of the sintered
ceramic materials. The viscosity of the ceramic slurry was
measured by a rotating rheometer (R/S+ Rheometer, Brookfield,
USA), and the test temperature was maintained at 25 ℃ using a
temperature controller (Brookfield-TC). A square (20 mm ×
20 mm) cured single layer was obtained by exposing the SiC slurry
inside a glass dish for 5–50 s using the DLP 3D printer digital
mask device (DMD). After cleaning and erasing the uncured
stock, a digital micrometer (211-101, Dongguan Sanliang
Measuring Tools Co., Ltd., China) was used to measure the curing
thickness for at least 5 points. The flexural strength was measured
in a microcomputer control electronic universal testing machine
(4505, Instron Experimental Equipment Trading Co., Ltd., USA)
54 J. Sun, J. Zhang, X. Zhang, et al.
J Adv Ceram2024,13(1):53−62
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using a three-point bending test (3PBT) with a loading speed of
0.5 mm/min and a span of 30 mm for at least 5 samples (3 mm ×
4 mm × 35 mm). Apparent porosity was determined by
Archimedes’ drainage hydrostatic method with at least 5 samples.
3Results and discussion
3.1Pre-oxidationofSiCpowders
The XRD patterns of the raw SiC powders and those after pre-
oxidation are shown in Fig. 1(a). The red line presented XRD of
the raw materials used in the experiment, which well matched the
standard diffraction pattern of 6H–SiC. After the pre-oxidation
treatment at 1200 ℃ for 2, 4, and 6 h, the crystal structure of the
SiC powders did not change significantly, but a new phase peak at
2θ of around 21.7°, which should be due to the formed cristobalite
(SiO2). A local magnification of the main diffraction peak of the
cristobalite phase is shown in Fig. 1(b). It can be seen that the
diffraction peak intensity of cristobalite gradually increased with
the prolongation of pre-oxidation time. The phase composition
change that occurred in the pre-oxidation process was that the SiC
powders oxidized in air at high temperatures generating
amorphous SiO2, which crystallized into cristobalite phases during
the cooling process as Reactions (1) and (2):
 +⇒() + 
()⇒()
The surface morphology of the raw SiC powders and that after
pre-oxidation were shown in Fig. 2. The original SiC powders
appeared as irregular shapes with relatively smooth surfaces as
shown in Fig. 2(a). After pre-oxidation, some tiny particles were
observed on the surface of the SiC particles, which might be due
to the chalking effect during pre-oxidation or ball milling
(Figs. 2(b)–2(d)). The surface of the SiC powders became rougher
with increasing the pre-oxidation time.
The element distribution of the SiC raw powders before
oxidation is shown in Figs. 3(a)–3(d), where C and Si were the
main elements on the surface. After the pre-oxidation treatment
(Figs. 3(e)–3(h)), besides C and Si, O was also detected on the
surface of the particles. The distribution of elements O and Si was
consistent with the profile of particles. Combined with the XRD
result in Fig. 1, it could be indicated that a scaly cristobalite layer
formed on the surface of the SiC powders after pre-oxidation
treatment.
The TEM image of pre-oxidized SiC powders is shown in
Fig. 4. A thin layer of SiO2 was found on the surface of the SiC
crystal, and its thickness was about 50–100 nm. The oxide shell
layer was non-uniform, which might be caused by the ball milling
process. Combined with the XRD results, the oxide layer was
cristobalite phase.
As particle size has a great effect on the viscosity of slurries [24],

Fig. 1  XRD patterns of raw SiC powders and those after pre-oxidation at 1200 ℃
for different time: (a) overall patterns; (b) local magnification of diffraction peak
at 21.7 ° in (a).

Fig. 2  SEM images of SiC powders before and after oxidation: (a) raw SiC
powders before oxidation; SiC powders oxidized at 1200 ℃ for (b) 2 h, (c) 4 h,
and (d) 6 h.

Fig. 3  SEM images and element distribution of SiC powders before and after oxidation: (a) raw SiC powders before oxidation; (e) SiC powders oxidized at 1200 ℃ for
4 h; (b–d) EDS mapping for (a) and (f–h) for (e).
High strength mullite-bond SiC porous ceramics fabricated by digital light processing 55
https://doi.org/10.26599/JAC.2024.9220835
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the particle size distribution of SiC powders before and after
oxidation at 1200 ℃ for different time was measured and shown
in Fig. 5. Pre-oxidized powders were ball milled and then sieved. It
can be seen that, d50 of raw SiC powders was 5.288 μm, which
changed to 4.851, 5.047, and 5.346 μm after oxidation at 1200 ℃
for 2, 4, and 6 h. The variation was quite small. Therefore, it can
be concluded that the viscosity of slurry made from SiC powders
before and after oxidation is not strongly related to particle size
changes.
3.2PhotocuringabilityofSiC-basedslurry
With experimental verification, it was found that the solid loading
of the pure fine SiC powder-based slurry was difficult to reach
35 vol% without additives. That was because the dispersion
behavior of the SiC powders inside the photosensitive resin was
poor, and the sedimentation of the SiC powders occurred quickly.
As a result, the SiC powder-based slurry without additives was not
suitable for printing. Herein, the raw SiC powders and pre-
oxidized ones were mixed with the sintering additives and added
to commercial resin to prepare the photocurable SiC slurry with a
solid content of 45 vol%. The curing ability of the slurry was now
discussed.
The curing layer thickness of the SiC ceramic slurry varied with
the exposure time of UV light as shown in Fig. 6. For the slurry
fabricated from raw SiC powders, no obvious curing/
polymerization occurred until the UV light exposure time was
above 20 s. With the increase in the exposure time, the curing
layer thickness increased. When the exposure time reached 45 s,
the curing layer thickness was around 46 μm. For the slurry with
pre-oxidized SiC powders, the curing performance was greatly

Fig. 4  TEM image of surface of SiC powders after oxidation: (a) TEM image of SiC powders; (b) high magnification image of circle region in (a); (c) inverse fast Fourier
transform pattern of area in (b); (d) inverse fast Fourier transform pattern of area in (b).

Raw materials
d10=1.820
d50=5.288
d90=8.656
1200 °C/2 h
d10=0.979
d50=4.851
d90=7.495
1200 °C/4 h
d10=0.899
d50=5.047
d90=8.849
1200 °C/6 h
d10=2.136
d50=5.346
d90=8.580
Fig. 5  Particle size distribution curve of SiC powders before and after oxidation: (a) raw SiC powdesr before oxidation; SiC powders oxidized at 1200 ℃ for (b) 2 h,
(c) 4 h, and (d) 6 h.
56 J. Sun, J. Zhang, X. Zhang, et al.
J Adv Ceram2024,13(1):53−62
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improved. Much shorter exposure time was required to obtain the
same curing thickness compared to the raw SiC powder-based
slurry. Among them, the ceramic slurry with the SiC powders
treated at 1200 ℃ for 4 h achieved the highest curing thickness at
all the range of the tested exposure time. To be specific, this slurry
exhibited a curing thickness as high as 51 μm with exposure time
of only 5 s, which was 40 s shorter than that of the control group.
As a result, the improved curing layer thickness can fully meet the
requirement for photocuring 3D printing, and therefore the
printing efficiency. According to the existing reports, the slurry
with un-treated SiC powders exhibited a curing thickness of < 30 μm
with exposure time of 90 s [24], whose exposure time is twice
while the curing thickness is thinner than that of the raw SiC
powder-based slurry in this work. In Ref. [27], 4.32 μm-sized SiC
powders were pre-oxidized at 1300 ℃, and the fabricated slurry
showed a curing thickness of 59 μm with exposure time of 5 s,
which was comparable with that of the treated SiC powder-based
slurry in this work.
With the extension of the oxidation time, the thickness of the
oxide layer on the SiC surface would increase, and the curing
ability of the ceramic slurry would be improved. This was
consistent with other reports [27].
The influence of the pre-oxidation time on curing properties of
the ceramic slurry exhibited a non-linear trend, i.e., at the same
exposure time, the curing layer thickness first increased with the
pre-oxidation time increasing from 2 to 4 h, and decreased with
the pre-oxidation time further increasing to 6 h. To explore the
influencing factors, the viscosity of the four ceramic slurry at the
solid loading of 45 vol% was tested as shown in Fig. 7(a). It can be
clearly seen that the pre-oxidation treatment effectively reduced
the viscosity of the SiC slurry. Among them, the slurry with the
1200 ℃/4 h treated SiC powders showed the highest viscosity in
the whole testing range of the shear rate. The viscosity values at
the shear rate of 30 s−1 are shown in Fig. 7(b). For the raw SiC
powder-based slurry, the viscosity was 2.543 Pa·s. For the slurry
made from the powders pre-oxidized at 1200 ℃ for 2, 4, and 6 h,
the viscosity decreased to 0.997, 1.369, and 0.859 Pa·s, respectively.

Fig. 6  Curve between curing thickness of SiC ceramic slurry with exposure time
for raw SiC powders and pre-oxidized SiC powders at 1200 ℃ for different
time.

Fig. 7  Viscosity of photosensitive SiC ceramic slurries: (a) viscosity at different shear rates and (b) viscosity at shear rate of 30 s−1 with different oxidation time and fixed
solid content of 45 vol%; (c) viscosity at different shear rates and (d) the viscosity at shear rate of 30 s−1 of slurry with different solid loadings prepared with SiC powders
pre-oxidized at 1200 ℃ for 4 h.
High strength mullite-bond SiC porous ceramics fabricated by digital light processing 57
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Effects of the solid loading on the viscosity of the ceramic slurry
fabricated from 4 h pre-oxidized SiC were further studied, and the
results are shown in Figs. 7(c) and 7(d). When the solid loading
increased from 45 to 50 vol%, the viscosity of the slurry slightly
increased from 1.369 to 3.112 Pa·s. With the solid loading further
increasing to 55 vol%, the viscosity of the slurry was still less than
10 Pa·s at a shear rate of 30 s−1, which is a criterion to judge
whether the slurry is suitable for printing [32]. To date, The
viscosity of the slurry with a pre-oxidized SiC particle size of
10 μm exceeded 10 Pa·s at a shear rate of 30 s−1 in a 55 vol% solid
content [26]. However, there have been few reports of the SiC
slurry with high solid loadings and low viscosity using powders
with a particle size of around 5 μm.
To investigate the effect of sintering additives on the viscosity,
three slurry were prepared separately as shown in Fig. 8. The
slurry A contained only pre-oxidized SiC powders. Based on this,
Al(OH)3 was added to form the slurry B. Also, in the slurry C,
along with pre-oxidized SiC powders, all additives, including
Al(OH)3, Y2O3, and CaF2 were added. When the solid load of the
slurry A reached 45 vol%, it showed a paste-like consistency with a
complete loss of flowability. Poor flowability was observed when
the solid content was 40 vol%. To facilitate the test, the solid
content of the slurry A was decreased to 35 vol%. The solid
content of 45 vol% was fabricated for the slurry B and C to
maintain the consistency with the above result.
As shown in Fig. 8(a), a strong shear-thinning behavior was
observed for the slurry A. The viscosity of both slurry B and C
almost became constant when the shearing rate was higher than
20 s−1. The viscosity of these slurry at a shear rate of 30 s−1 was
shown in Fig. 8(b). The visocisity of the slurry A was 1.375 Pa·s,
which changed to 2.232 and 1.369 Pa·s with the addition of
Al(OH)3 and Al(OH)3+Y2O3+CaF2, respectively. Considering the
lower solid content of the slurry A, we can conclude that the
addition of additives has a small effect on their viscosity.
Considering the best curing ability of the slurry made from the
1200 ℃/4 h treated SiC powders and also the acceptable viscosity
of the slurry at a solid loading of 50 vol%, this slurry was used to
print SiC bodies in the following work. To validate the feasibility
of using the slurry to print complex structures, a range of models
were designed, and printed, and sintered, as shown in Fig. 9.
Complex shapes were successfully maintained after sintering. The
result proved that the combination of pre-oxidation with the
addition of sintering additives was an effective way to enhance the
curing ability of the SiC slurry. The line shrinkage was similar for
the x- and y-axes in the horizontal direction with an average
shrinkage of 12.5%, and 22.7% for the z-axis along the print
direction. The total volume shrinkage was 41%.
The roughness of the top surface and side surface along the
print direction of the green body were measured and shown in
Fig. 10. The average surface roughness (Ra) of these two surfaces
was less than 1 μm, indicating the smooth feature of printing
surfaces.
The surface microstructure of the green body is shown in
Fig. 11. The top surface of the green body is distributed with
particles, and no obvious defects such as bumps and dents were
observed (Figs. 11(a) and 11(b)). In the side surface along the
printing direction, as shown in Figs. 11(c) and 11(d), printing
layers were observed without any other visible defects. The locally
enlarged image in Fig. 11(d) demonstrated that the combination
between layers was good, and there was no obvious layering
boundary and no debonding.
The top surface microstructure of the sintered body is shown in
Figs. 12(a) and 12(b). There were many pore structures distributed
on the top surface, and obvious rod-like grains were observed
around the pores. The particles were bonded together through the
glass phase, and no obvious crack defects were observed. On the
side surface along the printing direction, as shown in Figs. 12(c)
and 12(d), wave-like features were observed inheriting the layer-by-
layer printing characteristic. No interlayer debonding was noticed
from the locally enlarged image in Fig. 12(d).
3.3PropertiesofprintedSiCceramics
To avoid the negative impacts of the cristobalite phase generated
by pre-oxidation and oxidative debinding/sintering on the
performance of the SiC ceramics, Al(OH)3 was added as the Al2O3
source to react with SiO2 and form mullite, as illustrated in
Reactions (3) and (4):
()⇒+ 
+⇒·()
In addition, Y2O3 and CaF2 were added to enhance the
formation of mullite and the length/diameter ratio of the formed
mullite reinforcement. Based on the dry pressing process in our
previous work, the optimal amount of Al(OH)3 was 6.4 wt% of
the weight of SiC powders. However, the debinding treatment in
air was required for ceramics printed by the DLP method due to a
large amount of addition of organic matter such as photosensitive
resin and binders, which prolonged the entire oxidation process.
Besides, the SiC powders were also pre-oxidized before making

Fig. 8  Viscosity of photosensitive SiC ceramic slurry: (a) viscosity at different shear rates and (b) viscosity at shear rate of 30 s−1 with different compositions.
58 J. Sun, J. Zhang, X. Zhang, et al.
J Adv Ceram2024,13(1):53−62


the slurry. Therefore, it was necessary to explore and adjust the
proportion of the sintering additives to completely convert the
cristobalite phase into the mullite phase.
In this paper, four different amounts of Al(OH)3, i.e., 6.4, 10,
12.5, and 15 wt%, were added into samples, and their phase
composition after sintering was shown in Fig. 13. The crystalline
phase of the sintered samples included 6H–SiC, cristobalite, and
mullite as shown in Fig. 13(a). With the increase of the Al(OH)3
amount, the diffraction peak of the mullite phase was enhanced
and reached the highest when Al(OH)3 was 15 wt%, indicating
that some cristobalite continued to react with the Al2O3 source.
The peak intensity of the cristobalite phase first increased with the
Al(OH)3 amount increasing from 6.4 to 10 wt%, and then
disappeared at the Al(OH)3 amount of 12.5 wt%. Therefore, the
cristobalite produced in the whole process was completely
consumed at the Al(OH)3 amount of 12.5 wt%.
Interestingly, at the amount of Al(OH)3 to 15 wt%, a tiny
diffraction peak of the cristobalite phase reappeared. Since the
cristobalite phase has been eliminated at the Al(OH)3 addition
amount of 12.5 wt%, Al(OH)3 should be excessive. However, no
diffraction peaks of Al2O3 were detected in the XRD result,
indicating that the Al2O3 source had been consumed. Therefore, it
could be inferred that the excessive addition of Al(OH)3 might
accelerate the oxidation of SiC and therefore produce an extra
amount of the cristobalite phase [31,33], in which further
investigation is expected in the future.
Effects of the sintering temperature on the phase composition
of the SiC ceramic materials with 12.5 wt% Al(OH)3 addition were
shown in Fig. 13(b). With the increase of the sintering
temperature, the intensity of the diffraction peak of the cristobalite
phase decreased gradually, indicating that the higher sintering
temperature enhanced the entire mullitization process.
Theoretically, at a lower sintering temperature of 1450 ℃, the
content of cristobalite produced by oxidation during sintering
should be less than that at 1550 ℃, since a higher temperature
induces severer oxidation of SiC. However, as shown in Fig. 13(b),
a larger amount of the cristobalite phase was observed for the SiC
ceramic materials sintered at 1450 ℃ compared to that sintered at
1550 ℃. This was mainly because the sintering temperature of
1450 ℃ was not high enough to enable the mullite formation
reaction to completely conduct. As a result, increasing the
sintering temperature enhanced the formation reaction of mullite
and therefore the elimination of the cristobalite phase. As shown
in Fig. 13(b), with the increase of the sintering temperature, the
intensity of the diffraction peaks of mullite increased while that of
the cristobalite phase decreased. The cristobalite phase was out of
detection when the sintering temperature reached 1550 ℃. In
summary, the optimal addition amount of Al(OH)3 and sintering
temperature were found to be 12.5 wt% and 1550 ℃, respectively,
at which, oxidation-derived cristobalite was completely converted
into mullite phase, realizing mullite-bonded SiC ceramic
fabrication.
The microstructure and element distribution at the cross
section of the sintered SiC ceramic materials are shown in Fig. 14.
Although they were porous, the particles were bonded together,

Fig. 9  Photograph of complex structural SiC (a) green bodies and (b) sintered
bodies printed by DLP 3D printing.

Fig. 10  Surface roughness of green body: (a) 2D and (b) 3D image of top surface; (c) 2D and (d) 3D image of side surface along print direction where Ra is the average
roughness, Rq is the root-mean-square deviation, and Rmax is the maximum roughness depth.
High strength mullite-bond SiC porous ceramics fabricated by digital light processing 59
https://doi.org/10.26599/JAC.2024.9220835



Fig. 11  SEM images of printed green body: (a, b) top surface and (c, d) side surface in z-axis printing direction.

Fig. 12  SEM images of sintered ceramic body: (a, b) top surface and (c, d) side surface in z-axis printing direction.

Fig. 13  XRD patterns of SiC ceramics with different Al(OH)3 addition amounts and sintering temperatures: (a) different Al(OH)3 addition amounts sintered at
1550 ℃/2 h; (b) different sintering temperatures with 12.5 wt% Al(OH)3 addition.
60 J. Sun, J. Zhang, X. Zhang, et al.
J Adv Ceram2024,13(1):53−62


and no loose individual SiC particles were observed. According to
the EDS mapping result, Si, O, and Al elements were distributed
uniformly. This suggested that the mullite phase was formed as a
bonding phase between SiC particles making the materials well
sintered. The low sintering temperature of 1550 ℃ in air can also
effectively reduce the fabrication cost of complex SiC components
compared to the traditional high temperature/pressure and
vacuum/atmosphere sintering processes.
The flexural strength and apparent porosity of the SiC ceramic
materials sintered at 1550℃ in air by the DLP process were
shown in Fig. 15. The average flexural strength reached 97.6 MPa,
and the apparent porosity was 39.2 vol%. For comparison, the
performances of the SiC ceramic materials fabricated by other
techniques in literature, including the flexural strength and
porosity, were collected and listed in Table 1. The flexure strength
achieved in this study was 94.8% higher than that sintered at 2000 ℃
(50.1 MPa) in an Ar atmosphere by the DLP process, whose
porosity information remained unclear [27]. In addition, the
excellent flexural strength achieved by the DLP process in this
work was also comparable to that fabricated by other techniques,
e.g., SPS [34] and dry pressing method in our previous work [31]
with similar porosity, and much higher than the SiC ceramic
materials fabricated by dry pressing and sintering from other
groups.
The flexural strength of SiC porous ceramics prepared in this
study was high compared to the sintered bodies with similar
apparent porosity, but high porosity was a limitation to the
ceramics’ properties. The density of sintered ceramics could be
improved by subsequent post-treatment processes, such as PIP,
CVI, and RMI techniques, to enhance their properties [8,35−38].
Ding et al. [8] and Chen et al. [37] greatly enhanced the density
and strength of SiC and C/SiC by PIP post-treatment [8,37].
Other researchers combined PIP with CVI [36,38] or CVI with
RMI [35], and the properties of SiC-based composites were greatly
enhanced. These treatments provided practical ideas for our
future research.
4Conclusions
In this paper, complex-shaped SiC ceramic materials with high
flexural strength were successfully fabricated by the DLP
technique and low-temperature and air-atmosphere sintering.
With the pre-oxidation treatment, SiO2 shell layers formed on the
SiC powders. As a result, the printing ability of the SiC
photosensitive slurry was effectively enhanced even though the
SiC particle size was as fine as 5 μm: a high curing thickness was
achieved above 50 μm with exposure time as short as 5 s; the
viscosity of the SiC slurry with a solid loading of 55 vol% was as
low as 7.77 Pa·s at a shear rate of 30 s−1. Moreover, with the
rational design of the sintering additives, the oxidation-derived
cristobalite in the materials was completely consumed and

Fig. 14  Micromorphology and element distribution of cross section of sintered SiC ceramic materials.

Fig. 15  Flexural strength and apparent porosity of SiC ceramic materials
sintered at 1550 ℃ in air by DLP technique.

Table 1  Comparison of apparent porosity and flexural strength of SiC porous ceramics prepared by photocuring and traditional methods
Material; additive Process Sintering condition Flexural strength (MPa) Apparent porosity (vol%) Ref.
SiC; Al(OH)3+Y2O3+CaF2DLP Pressureless, 1550 ℃, air 97.6 39.2 This work
SiC DLP Pressureless, 2000 ℃, argon 47.9 — [27]
SiC; Y2O3+Al2O3SLA Pressureless, 1950 ℃, nitrogen 229.0 10.2 [39]
SiC — SPS, 1800 ℃, vacuum 103.0 35.7 [34]
SiC; MoO3+Al2O3Dry pressing Pressureless, 1000 ℃, air 66.0 45.4 [40]
SiC; Al(OH)3+Y2O3+CaF2Dry pressing Pressureless, 1550 ℃, air 113.0 40.3 [31]
SiC; fly ash+MoO3Dry pressing Pressureless, 1000 ℃, air 38.4 36.4 [41]
SiC; Al2O3+graphite Dry pressing Pressureless, 1450 ℃, air 27.5 44.4 [42]
High strength mullite-bond SiC porous ceramics fabricated by digital light processing 61
https://doi.org/10.26599/JAC.2024.9220835


converted into a higher-strength mullite bonding phase. The thus
fabricated SiC ceramic materials exhibited flexural strength of
97.6 MPa at apparent porosity of 39.2 vol%, which is comparable
to that fabricated by the SPS process. This work therefore provides
a novel method to fabricate DLP-process SiC ceramic materials
with excellent printing ability and high mechanical strength.
Acknowledgements
This work was supported by Shandong University−MSEA
International Institute for Materials Genome Joint Innovation
Center for Advanced Ceramics, and the Key Research and
Development Projects of Shaanxi Province (Nos. 2018ZDCXL-
GY-09-06 and 2021ZDLGY14-06).
Declarationofcompetinginterest
The authors have no competing interests to declare that are
relevant to the content of this article.
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