Influence of pre-synthesized Al2O3-SiC composite powder from clay on properties of low-carbon MgO-C refractories

Abstract

To improve the properties of low-carbonization of MgO–C refractories, the introduction of composite additives is an effective strategy. Al2O3–SiC composite powder was prepared from clay using electromagnetic induction heating and carbon embedded methods. Further, the Al2O3–SiC composite powder synthesized by electromagnetic induction heating at 600 A was added into low-carbon MgO–C refractories (4 wt.%) to improve their properties. The results showed that when the addition amount of Al2O3–SiC composite powder is within the range of 2.5–5.0 wt.%, the properties of low-carbon MgO–C samples were significantly improved, e.g., the apparent porosity of 7.58%–8.04%, the bulk density of 2.98–2.99 g cm−3, the cold compressive strength of 55.72–57.93 MPa, the residual strength after three air quenching at 1100 °C of 74.86%–78.04%, and the decarburized layer depth after oxidized at 1400 °C for 2 h of 14.03–14.87 mm. Consequently, the idea for the rapid synthesis of Al2O3–SiC composite powder provides an alternative low-carbon MgO–C refractories performance optimization strategy.

Full text

ORIGINAL PAPER
Influence of pre-synthesized Al
2
O
3
–SiC composite powder from clay
on properties of low-carbon MgO–C refractories
Bei-yue Ma
1,2
•Xin-ming Ren
1,2
•Zhi Gao
1,2
•Fan Qian
3
•Zhao-yang Liu
2
•Guo-qi Liu
3
•Jing-kun Yu
2
•
Gao-feng Fu
1,2
Received: 27 April 2021 / Revised: 2 July 2021 / Accepted: 26 July 2021
China Iron and Steel Research Institute Group 2021
Abstract
To improve the properties of low-carbonization of MgO–C refractories, the introduction of composite additives is an
effective strategy. Al
2
O
3
–SiC composite powder was prepared from clay using electromagnetic induction heating and
carbon embedded methods. Further, the Al
2
O
3
–SiC composite powder synthesized by electromagnetic induction heating at
600 A was added into low-carbon MgO–C refractories (4 wt.%) to improve their properties. The results showed that when
the addition amount of Al
2
O
3
–SiC composite powder is within the range of 2.5–5.0 wt.%, the properties of low-carbon
MgO–C samples were significantly improved, e.g., the apparent porosity of 7.58%–8.04%, the bulk density of
2.98–2.99 g cm
-3
, the cold compressive strength of 55.72–57.93 MPa, the residual strength after three air quenching at
1100 C of 74.86%–78.04%, and the decarburized layer depth after oxidized at 1400 C for 2 h of 14.03–14.87 mm.
Consequently, the idea for the rapid synthesis of Al
2
O
3
–SiC composite powder provides an alternative low-carbon MgO–C
refractories performance optimization strategy.
Keywords MgO–C refractory Electromagnetic induction heating Al
2
O
3
–SiC composite powder Oxidation resistance
1 Introduction
MgO–C refractories, a kind of the largest market demand
carbon-containing refractories, have been widely used in
metallurgical kiln linings, such as alternating current/direct
current electric arc furnace, converter, Ruhrstal Heraeus
refining furnace, and ladle furnace slag line, because of
excellent high-temperature performance and outstanding
basic slag resistance [1–5]. The rapid development of
special steels, including ultra-low-carbon steel, has further
promoted the low-carbonization of MgO–C refractories.
On the one hand, refractories with high carbon content may
cause molten steel to entrap slag inclusions, resulting in
carbon content out of standard. On the other hand, although
graphite with high thermal conductivity brings better
thermal shock resistance to refractories, it reduces strength
and increases energy consumption of per ton steel [6–8].
Therefore, the low-carbonization of MgO–C refractories,
as well as the strengthening and toughening after low-
carbonization, is the research focuses currently in the field
of refractory materials.
Due to the quantum size effect, surface effect, and high
dispersibility of nanocarbon, MgO–C refractories using
nanocarbon as a carbon source can form a complete and
continuous carbon network inside even with very low-
carbon content; thus, they still exhibit high performance in-
line with expectation [9]. Liu et al. [10] investigated the
effects of different types of carbon black on the properties
of low-carbon MgO–C refractories (3 wt.% C) and found
that the low-carbon MgO–C refractories added with
0.4 wt.% N220 carbon black showed the almost identical
thermal shock resistance better as traditional MgO–C
refractories (16 wt.% C). Zhu et al. [11] found that the
&Xin-ming Ren
maceramic@126.com
1
Key Laboratory for Ecological Metallurgy of Multimetallic
Mineral (Ministry of Education), Shenyang 110819,
Liaoning, China
2
School of Metallurgy, Northeastern University,
Shenyang 110819, Liaoning, China
3
State Key Laboratory of Advance Refractories, Sinosteel
Luoyang Institute of Refractories Research Co., Ltd.,
Luoyang 471039, Henan, China
123
J. Iron Steel Res. Int.
https://doi.org/10.1007/s42243-021-00653-8(0123456789().,-volV)(0123456789().,-volV)

introduction of graphite oxide nanosheets and carbon
nanotubes enhanced the cold compressive strength and
thermal shock resistance of low-carbon MgO–C refracto-
ries (5 wt.% C), respectively. However, although nanos-
tructured carbon makes up for the shortcomings of thermo-
mechanical properties for low-carbon MgO–C refractories,
it brings greater challenges in terms of oxidation resistance.
Over the past decade, a series of high-quality antioxidants
have been reported, including some single-phase powders
with higher oxidation resistance like Al–Mg [12], Cr
7
C
3
[6], Ti
3
AlC
2
[8], Al
4
SiC
4
[13], Al
4
O
4
C[14], and MgB
2
[15], as well as composite powder with both antioxidation
and toughening functions like B
4
C/Al
2
O
3
/C [16], Al
2
O
3
/
SiC [17], C/MgAl
2
O
4
[18], and TiC/Ti
3
AlC/C [19]. In
contrast, composite powder has more development poten-
tial because of easier dispersion and lower cost. Among
them, Al
2
O
3
–SiC composite powder shows the best
research potential due to its wide range of sources and
efficient application results.
For this purpose, the present work used abundant low-
cost clay as the main raw material to prepare Al
2
O
3
–SiC
composite powder by electromagnetic induction heating
method and conventional carbon embedded method. Based
on the relevant thermodynamic data, the synthesis mech-
anism of Al
2
O
3
–SiC composite powder was analyzed.
Furthermore, the as-prepared Al
2
O
3
–SiC composite pow-
der was used to replace the MgO fine powders in low-
carbon MgO–C refractories, and the effects of the com-
posite powder addition amounts on their thermal shock
resistance and oxidation resistance were investigated in
detail.
2 Experimental
2.1 Raw materials
As-received refractory clay (from Jilin, China) and carbon
black (N990 type; B300 nm; from Qingdao, China) were
used as the starting materials for synthesizing Al
2
O
3
–SiC
composite powder. Table 1and Fig. 1show the X-ray
fluorescence results, X-ray diffraction (XRD) patterns, and
particle size distribution of clay, respectively.
Fused magnesia (particle sizes of B0.088 mm,
0–1 mm, 1–3 mm, and 3–5 mm; MgO purity C97 wt.%;
from Yingkou, China), natural flake graphite
(B0.088 mm; fixed carbon purity C97.5 wt.%; from
Yingkou, China), aluminum powder (B0.088 mm; Al
purity C98 wt.%; from Yingkou, China), and liquid phe-
nolic resin (carbon yield C43.0 wt.%; from Yingkou,
China) were used as the raw materials for preparing the
low-carbon MgO–C refractories.
2.2 Sample preparation
The synthesis process of Al
2
O
3
–SiC composite powder is
as follows. Clay and carbon black were quantitatively
weighed according to the mass ratio of 10:4, and then the
raw powders were poured into a high-energy planetary ball
mill equipped with cemented carbide jars (QM-3SP4,
China) and ball milled at the rate of 200 r min
–1
for 4 h.
The homogeneous raw powders were pressed into cylin-
drical green bodies (/20 mm 920 mm) at 20 MPa.
Finally, the green bodies were heated in a high-temperature
box resistance furnace (in carbon embedded atmosphere at
1500, 1550, and 1600 C for 4 h) and an electromagnetic
induction heating furnace (in Ar atmosphere of 0.3 L min
–1
at 400, 500 and 600 A for 15 min), respectively.
The preparation process of low-carbon MgO–C refrac-
tories is as follows. The raw materials were weighed
according to the formulas shown in Table 2, and then
mixed progressively in a roll-type mixer with addition of
fused magnesia aggregate, liquid phenolic resin, flake
graphite, magnesia powder, Al powder, and Al
2
O
3
–SiC
composite powder in sequence. The homogeneous raw
materials were pressed into different batches of cylindrical
green bodies (/50 mm 950 mm) at 200 MPa. Finally, the
green bodies were cured at 200 C for 24 h and then coked
at 1400 C for 3 h.
2.3 Characterization
In this work, the formed phases, microstructure, and the
elements in the local microstructure of the samples were
detected using an X-ray diffractometer (Rigaku-Ultima IV,
Japan; 5 ()min
–1
,10–80), a scanning electron micro-
scope (SEM, Hitachi-S4800, Japan), and an energy dis-
persive spectrometer (EDS), respectively. The apparent
porosity, bulk density, cold compressive strength, and
thermal shock resistance of the samples were measured
based on the corresponding Chinese standards (GB/T
2997–2000, 5072–2008, 30873–2014, and 21650.1–2008).
Specifically, the thermal shock resistance of the sintered
samples was tested by air quenching method and quanti-
tatively characterized by the residual cold compressive
strength of the samples after air quenching three times at
1100 C. In addition, the oxidation test of the samples was
carried out at 1400 C for 2 h [20]. For all the above tests,
Table 1 Chemical compositions of clay (wt.%)
SiO
2
Al
2
O
3
Fe
2
O
3
TiO
2
K
2
O MgO Others
63.26 32.34 1.68 1.16 0.90 0.28 0.38
B.Y. Ma et al.
123

the average value of the recorded results of three parallel
samples was selected as the final test results.
3 Results and discussion
3.1 Synthesis of Al
2
O
3
–SiC composite powder
from clay by two methods
3.1.1 Phase compositions
Figure 2shows XRD patterns of Al
2
O
3
–SiC composite
powder synthesized by electromagnetic induction heating
method at different current densities (Fig. 2a–c) and carbon
embedded method at different temperatures (Fig. 2d–f).
From Fig. 2a–c, the phases of the samples are mullite and
cristobalite at the current intensity of 400 A (Fig. 2a); the
phase of the samples are mullite, corundum, and silicon
carbide (b-type) at the current intensity of 500 A (Fig. 2b);
the phase of the samples are corundum and silicon carbide
at the current intensity of 600 A (Fig. 2c). Similarly, with
the temperature increasing, the phases of the samples
transform from mullite and silicon carbide (1500 C,
Fig. 2d) to mullite, corundum, and silicon carbide
(1550 C, Fig. 2e) and then to silicon carbide and corun-
dum (1600 C, Fig. 2f). The XRD results suggest that
increasing the current intensity or rising the temperature is
beneficial to the formation of target phase of Al
2
O
3
and
SiC, and when the temperature is 1600 C or the current
intensity is 600 A, the content of the impurity phases in the
samples is the lowest.
3.1.2 Microstructure evolution
Figure 3presents the SEM images and EDS results of
Al
2
O
3
–SiC composite powder synthesized by electromag-
netic induction heating method at 600 A (Fig. 3a–c) and
carbon embedded method at 1600 C (Fig. 3d–f). The
powder samples synthesized by electromagnetic induction
heating method are composed of two kinds of particles.
Combined with the EDS results, the large particles with a
size of 5–10 lm are alumina, and the small particles with a
size of 10–20 lm are silicon carbide. Likewise, the powder
samples synthesized by carbon embedded method are also
composed of large particles of alumina and small particles
of silicon carbide. By contrast, the particle size distribution
of the powder samples synthesized by electromagnetic
induction heating is more uniform than that of the powder
Fig. 1 XRD pattern (a) and particle size distribution (b) of clay
Table 2 Formulas of low-carbon MgO–C refractories
Material Particle size/mm MC/wt.% MC-2.5AS/wt.% MC-5AS/wt.% MC-7.5AS/wt.%
Fused magnesia aggregate (0–1) ?(1–3) ?(3–5) 77.0 77.0 77.0 77.0
Magnesia powder B0.088 17.0 14.5 12.0 9.5
Flake graphite B0.088 4.0 4.0 4.0 4.0
Al powder B0.088 2.0 2.0 2.0 2.0
Al
2
O
3
–SiC composite powder – 0.0 2.5 5.0 7.5
Liquid phenolic resin – ?4.0 ?4.0 ?4.0 ?4.0
Influence of pre-synthesized Al
2
O
3
–SiC composite powder from clay on properties of low-carbon…
123

samples synthesized by carbon embedded method. The
results fully reflect the difference between the two techni-
cal methods. The heating rate of electromagnetic induction
heating is higher, and the preset temperature can be
reached in a very short time [21]. When electromagnetic
induction heating was used to treat the clay, the mullite
grains inside powder samples were reduced to alumina and
silicon carbide by carbon black before they grew up.
Naturally, the particle size of reduction products, especially
alumina particles, is smaller than that of the samples syn-
thesized by carbon embedded method.
3.1.3 Analysis of synthesis process of Al
2
O
3
–SiC composite
powder
From the XRD results in Sect. 3.1.1, when electromagnetic
induction heating is used at the current of 400 A
(*1300 C), the phases of the powder samples are
cristobalite and mullite; as the current increases to 500 A or
the temperature rises to 1550 C (carbon embedded
method), the diffraction peaks of cristobalite disappeared
while the diffraction peaks of silicon carbide and corundum
emerged; when the current increases to 600 A or the
temperature rises to 1600 C, all mullite was transformed
into silicon carbide and corundum. Therefore, the reduction
process of clay can be more specifically divided into three
successive stages: decomposition of kaolinite (reactions
(1)–(4)), reduction of cristobalite (reactions (5)–(7)), and
reduction of mullite (reactions (8) and (9)) [22].
Al2O32SiO22H2OsðÞ¼ Al2O32SiO2sðÞþ
2H2OgðÞ 600 CðÞ
ð1Þ
3Al
2O32SiO2
ðÞsðÞ¼ Al6Si2O13 sðÞþ
4SiO2amorphousðÞ950 CðÞ
ð2Þ
Al6Si2O13 sðÞ¼ 3Al2O32SiO2sðÞ 1250 CðÞð3Þ
SiO2amorphousðÞ¼SiO2ðsÞ1200 CðÞ ð4Þ
CsðÞþ SiO2sðÞ¼ SiO gðÞþCO gðÞ ð5Þ
2C sðÞþ SiO gðÞ¼SiC sðÞþ CO gðÞ ð6Þ
3C sðÞþ SiO2sðÞ¼ SiC sðÞþ 2CO gðÞ ð7Þ
Fig. 2 XRD patterns of Al
2
O
3
–SiC composite powder synthesized by electromagnetic induction heating method (a–c) and carbon embedded
method (d–f). a400 A; b500 A; c600 A; d1500 C; e1550 C; f1600 C
B.Y. Ma et al.
123

3Al2O32SiO2sðÞþ 4SiO2sðÞþ 18C sðÞ¼
3Al2O3sðÞþ 6SiC sðÞþ 12CO gðÞ ð8Þ
3Al2O32SiO2sðÞþ 6C sðÞ¼ 3Al2O3sðÞþ 2SiC sðÞþ
4CO gðÞ
ð9Þ
Among them, reactions (5)–(9) involve gas phases; thus,
when calculating their Gibbs free energy values, it is
necessary to take into account the influence of the gas
phases partial pressure in the real reaction system. Taking
reaction (7) for instance, its Gibbs free energy equals
DG7¼DGH
7þRT ln J¼16:63Tln PCO=PH331:98Tþ
603;150, where DGHis the standard Gibbs free energy, J
mol
-1
;Ris the molar gas constant, 8.314 J mol
-1
J
-1
;Tis
the initial temperature, K; Jis the pressure quotient; P
CO
is
the partial pressure of CO, Pa; and P
H
is the standard
atmospheric pressure, 1 910
5
Pa. The calculation results
show that in standard atmosphere, the initial temperature of
the spontaneous reaction is 1543.68 C; in Ar atmosphere
(P
CO
= 20 Pa), the initial temperature is reduced to
1000.34 C; in carbon embedded atmosphere
(P
CO
= 42,000 Pa), the initial temperature is decreased to
1452.22 C. By comparison, the initial temperature of the
reduction reaction in Ar atmosphere is lower, indicating
that the electromagnetic induction method is a better
choice to synthesize Al
2
O
3
–SiC composite powder.
Therefore, the overall reaction of the clay reduction
process is as follows [23]:
Al2O32SiO22H2OsðÞþ 6C sðÞ¼
Al2O3sðÞþ 2SiC sðÞþ 4CO gðÞþ2H2OgðÞ ð10Þ
3.2 Effects of as-synthesized Al
2
O
3
–SiC
composite powder on properties of low-
carbon MgO–C refractories
3.2.1 Phase compositions and microstructures
Figure 4a shows the XRD patterns of samples MC and
MC-5AS after coked at 1400 C for 3 h. The diffraction
peaks of graphite and periclase phases were found in the
XRD pattern of sample MC. In comparison, in addition to
Fig. 3 SEM images and EDS results of Al
2
O
3
–SiC composite powder synthesized by electromagnetic induction heating method at 600 A (a–
c) and carbon embedded method at 1600 C(d–f)
Influence of pre-synthesized Al
2
O
3
–SiC composite powder from clay on properties of low-carbon…
123

graphite and periclase phases, the diffraction peaks of
magnesia-alumina spinel and forsterite phases were also
detected in sample MC-5AS. In order to further investigate
the influence of Al
2
O
3
–SiC composite powder on the
microstructure of the low-carbon MgO–C samples, the
SEM images of the samples are displayed in Fig. 4b–e.
MgO aggregates (light grey) and flake graphite (dark grey)
together form the matrix structure of the samples, and some
pores (black) are dispersed between the particles. With the
increase in the addition amount of Al
2
O
3
–SiC composite
powder, the number of pores on the surface of the samples
(from sample MC to sample MC-5AS) decreases, and the
pore size becomes smaller. However, bigger and more
pores reappear on the surface of sample MC-7.5AS,
especially where the spinel and forsterite phases are enri-
ched (confirmed by EDS).
3.2.2 Physical properties
Figure 5shows the apparent porosity and bulk density of
the MgO–C samples after coked at 1400 C for 3 h. The
apparent porosity of the MgO–C samples first decreases
and then increases with the increase in the Al
2
O
3
–SiC
composite powder addition, and the sample MC-2.5AS
achieve the minimum apparent porosity of 7.58%. In
contrast, the bulk density of the MgO–C sample rises first
and then decreases, and the corresponding maximum value
of the bulk density is obtained at the 2.5 wt.% point. This
change can be mainly attributed to two reasons. One is that
the micron-sized Al
2
O
3
–SiC composite powder has good
dispersibility and can be better filled between the aggre-
gates, thereby bringing the microstructure with high
packing to the MgO–C samples [24]. The other is that the
Al
2
O
3
–SiC composite powder can react with MgO in the
MgO–C sample to form spinel. According to theoretical
density calculations, *8% volume expansion will gener-
ate when Al
2
O
3
and MgO react to form stoichiometric
spinel, which is beneficial to offsetting some pore defects
inside the material [25]. However, when the volume
expansion caused by the solid–solid reaction is too large,
Fig. 4 XRD patterns (a) and SEM images (b–e) of MgO–C samples. bSample MC; csample MC-2.5AS; dsample MC-5AS; esample MC-
7.5AS. FM Fused magnesia; FG flake graphite; MA spinel; M
2
S forsterite; P pore
Fig. 5 Al
2
O
3
–SiC composite powder addition amounts versus appar-
ent porosity and bulk density of samples
B.Y. Ma et al.
123

the original structure of the material will inevitably be
destroyed, resulting in the formation of some new cracks
and pores. This is also the reason why the apparent porosity
of samples MC-5AS and MC-7.5AS is slightly higher than
that of sample MC-2.5AS (but still better than sample MC).
3.2.3 Cold compressive strength and thermal shock
resistance
Figure 6shows the cold compressive strength of the MgO–
C samples before and after quenched at 1100 C for 3
times and the corresponding residual strength ratio. The
strength of the MgO–C samples, as expected, displays a
trend of first increasing and then decreasing, same as the
apparent porosity of the samples. The change in the cold
compressive strength of the MgO–C samples can be
explained by their microstructures. With the introduction of
Al
2
O
3
–SiC composite powder, the aggregate and matrix
(fine powders) of the MgO–C samples achieve a more ideal
close packed structure, and meanwhile, the in-situ formed
ceramic phases (spinel and forsterite) further enhance the
bond between the aggregates (Fig. 4b–e). As shown in
Fig. 6, the residual strength ratio of the MgO–C samples
increases from 67.86% (MC) to 74.86% (MC-2.5AS),
78.04% (MC-5AS), and 71.47% (MC-7.5AS), respectively.
This can be attributed to the effect of microcracks and
micropores formed due to volume expansion. For ceramics
and refractories, generally, although the generation of
microcracks and micropores will inevitably increase the
apparent porosity of the material to a certain extent, they
can even paly a toughening effect to improve the high-
temperature properties of the material if controlled prop-
erly. Therefore, for sample MC-5AS, although the micro-
cracks and micropores caused the increase in its apparent
porosity, these microcracks and pores brought it better
high-temperature properties including thermal shock
resistance. In addition, for refractory materials, thermal
shock damage and fracture are dominated by crack prop-
agation rather than generation. Therefore, the pre-forming
microcracks and micropores inside samples MC-2.5AS and
MC-5AS effectively delay the propagation of cracks by
increasing the surface energy. Conversely, the size of the
cracks inside sample MC-7.5AS exceeds the critical
stable crack length, which causes the crack propagation
process to be accelerated [26,27].
3.2.4 Oxidation resistance
Figure 7a and b presents the optical images, the corre-
sponding oxidation ratio and decarburized layer depth of
the MgO–C samples after oxidized at 1400 C for 2 h.
With the increase in the addition amount of Al
2
O
3
–SiC
composite powder, the decarburized layer area (light color)
of the MgO–C samples decreases first and then increases.
The results calculated by image analysis software indicate
that sample MC-5A has the best oxidation resistance,
which is represented by the lowest oxidation ratio and
thickness of the decarburized layer. It is generally believed
that the oxidation resistance of refractories is directly
related to the apparent porosity, because the open pores
provide the main channel for the infiltration of oxygen
[28]. The apparent porosity of samples MC-2.5AS and
MC-5AS is lower than that of samples MC and MC-7.5AS;
thus, there are fewer channels for oxygen to enter them. In
addition, the chemical reaction that occurs in the decar-
burized layer will also affect the subsequent oxidation. The
XRD pattern of sample MC-5AS after oxidation is shown
in Fig. 7c. Except for MgO phase, both MgAl
2
O
4
and
Mg
2
SiO
4
phases are newly formed oxidation products, and
their reaction formula is as follows:
Al2O3sðÞþ MgO sðÞ¼ MgAl2O4sðÞ ð11Þ
SiC sðÞþ 2MgO sðÞþ O2gðÞ¼Mg2SiO4sðÞþ CsðÞ
ð12Þ
On the one hand, the volume expansion of reactions (11)
and (12) filled the pores of the oxide layer and formed a
dense layer in the oxide area of sample MC-5AS, thereby
reducing the possibility of continued contact between
oxygen and graphite. On the other hand, reaction (12) made
up for the loss of carbon inside sample MC-5AS, thereby
further delaying the oxidation process.
Fig. 6 Cold compressive strength and thermal shock resistance of
samples versus Al
2
O
3
–SiC composite powder addition amounts
versus
Influence of pre-synthesized Al
2
O
3
–SiC composite powder from clay on properties of low-carbon…
123

4 Conclusions
1. As the current intensity increased, the phases of clay
transformed from mullite and cristobalite (400 A) to
mullite, corundum, b-type silicon carbide (500 A) and
then to corundum and silicon carbide (600 A); simi-
larly, when the carbon embedded was used, silicon
carbide was detected at 1500 C, and the phases of the
samples were transformed into the target phases of
Al
2
O
3
and SiC at 1600 C. In comparison, the elec-
tromagnetic induction heating method can save more
time and energy, and meanwhile, the particle size
distribution and morphology of the Al
2
O
3
–SiC com-
posite powder synthesized by electromagnetic induc-
tion heating method are superior to those of the Al
2
O
3
–
SiC composite powder synthesized by carbon embed-
ded method.
2. When 2.5 wt.% Al
2
O
3
–SiC composite powder is
added, the sintering properties of low-carbon MgO–C
samples are the best, which can be attributed to the
volume expansion effect generated by the reaction
between Al
2
O
3
and MgO matrix. When 5 wt.%
Al
2
O
3
–SiC composite powder is introduced, the
microcrack toughening effect makes sample MC-5AS
exhibit the best thermal shock resistance. In addition,
Al
2
O
3
–SiC composite powder improves the oxidation
resistance of low-carbon MgO–C samples by forming
a spinel- and forsterite-containing dense protective
layer in the oxide layer.
Acknowledgements This work was financially supported by the
National Natural Science Foundation of China (Grant Nos.
U20A20239 and U1908227), the Fundamental Research Funds for the
Central Universities (Grant No. N2125002), and the open research
fund for State Key Laboratory of Advance Refractories (Grant No.
SKLAR202001).
Fig. 7 Optical images of cross-section (a), oxidation ratio and decarburized layer depth (b) of samples after oxidized at 1400 C for 2 h, and
XRD pattern of decarburized layer of sample MC-5AS (c)
B.Y. Ma et al.
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References
[1] Y. Chen, C.J. Deng, X. Wang, J. Ding, C. Yu, H.X. Zhu, Constr.
Build. Mater. 240 (2020) 117964.
[2] C. Wo
¨hrmeyer, S. Gao, Z.F. Ping, C. Parr, C.G. Aneziris, P.
Gehre, Steel Res. Int. 91 (2020) 1900436.
[3] S.E. Gass, P.G. Galliano, A.G. Tomba Martinez, J. Eur. Ceram.
Soc. 41 (2021) 3769–3781.
[4] M. Chen, S. Gao, L. Xu, N. Wang, Ceram. Int. 45 (2019)
21023–21028.
[5] X.M. Ren, B.Y. Ma, S.M. Li, H.X. Li, G.Q. Liu, W.G. Yang, F.
Qian, S.X. Zhao, J.K. Yu, J. Iron Steel Res. Int. 28 (2021)
38–45.
[6] Y. Yang, J. Yu, H.Z. Zhao, H. Zhang, P.D. Zhao, Y.C. Li, X.Q.
Wang, G.P. Li, Ceram. Int. 46 (2020) 19743–19751.
[7] T.B. Zhu, Y.W. Li, S.B. Sang, J. Alloy. Compd. 783 (2019)
990–1000.
[8] J.F. Chen, N. Li, J. Huba
´lkova
´, C.G. Aneziris, J. Eur. Ceram.
Soc. 38 (2018) 3387–3394.
[9] S. Behera, R. Sarkar, Prot. Met. Phys. Chem. Surf. 52 (2016)
467–474.
[10] B. Liu, J.L. Sun, G.S. Tang, K.Q. Liu, L. Li, Y.F. Liu, J. Iron
Steel Res. Int. 17 (2010) No. 10, 75–78.
[11] T.B. Zhu, Y.W. Li, S.B. Sang, S.L. Jin, Y.B. Li, L. Zhao, X.
Liang, Ceram. Int. 40 (2014) 4333–4340.
[12] A.P. Luz, T.M. Souza, C. Pagliosa, M.A.M. Brito, V.C. Pan-
dolfelli, Ceram. Int. 42 (2016) 9836–9843.
[13] H.B. Yao, X.M. Xing, E.H. Wang, B. Li, J.H. Chen, J.L. Sun,
X.M. Hou, Coatings 7 (2017) 85.
[14] C. Yu, H.X. Zhu, W.J. Yuan, C.J. Deng, P. Cui, S.M. Zhou, J.
Alloy. Compd. 579 (2013) 348–354.
[15] K.S. Campos, G.F.B. Lenz e Silva, E.H.M. Nunes, W.L. Vas-
concelos, Ceram. Int. 38 (2012) 5661–5667.
[16] D.H. Ding, X.C. Chong, G.Q. Xiao, L.H. Lv, C.K. Lei, J.Y. Luo,
Y.F. Zang, Ceram. Int. 45 (2019) 16433–16441.
[17] B.Y. Ma, Q. Zhu, Y. Sun, J.K. Yu, Y. Li, J. Mater. Sci. Technol.
26 (2010) 715–720.
[18] L.H. Lv, G.Q. Xiao, D.H. Ding, Y. Ren, S.L. Yang, P. Yang, X.
Hou, Int. J. Appl. Ceram. Technol. 16 (2019) 1253–1263.
[19] X.G. Liu, S.W. Zhang, Y. Li, X.Y. Yang, Y. Zhang, Ceram. Int.
44 (2018) 22567–22573.
[20] X.M. Ren, B.Y. Ma, G.L. Zhang, G.F. Fu, J.K. Yu, G.Q. Liu,
Mater. Chem. Phys. 252 (2020) 123309.
[21] J. Chen, D. Chen, H.Z. Gu, A. Huang, H.W. Ni, J. Alloy.
Compd. 870 (2021) 159463.
[22] B.Y. Ma, X.M. Ren, Y. Yin, L. Yuan, Z. Zhang, Z.Q. Li, G.Q.
Li, Q. Zhu, J.K. Yu, Ceram. Int. 43 (2017) 11830–11837.
[23] A.C.D. Chaklader, S.D. Gupta, E.C.Y. Lin, B. Gutowski, J. Am.
Ceram. Soc. 75 (1992) 2283–2285.
[24] D.H. Ding, L.H. Lv, G.Q. Xiao, J.Y. Luo, C.K. Lei, Y. Ren, S.L.
Yang, P. Yang, X. Hou, Int. J. Appl. Ceram. Technol. 17 (2020)
645–656.
[25] M.T. Li, N.N. Zhou, X.D. Luo, G.D. Zhang, Z.P. Xie, L.C. Xu,
P.C. Liu, Mater. Chem. Phys. 175 (2016) 6–12.
[26] I. Khlifi, O. Pop, J.C. Dupre
´, P. Doumalin, M. Huger, J. Eur.
Ceram. Soc. 39 (2019) 3893–3902.
[27] Y.J. Dai, Y.W. Li, X.F. Xu, Q.Y. Zhu, W. Yan, S.L. Jin, H.
Harmuth, J. Eur. Ceram. Soc. 39 (2019) 5433–5441.
[28] Y.M. Qin, X.G. Liu, Q. Zhang, F. Zhao, X.H. Liu, Q.L. Jia,
Corros. Sci. 166 (2020) 108446.
Influence of pre-synthesized Al
2
O
3
–SiC composite powder from clay on properties of low-carbon…
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