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Low carbon containing Al2O3-C refractories with nanocarbon as the sole carbon source
Journal Pre-proof Low carbon containing Al2O3-C refractories with nanocarbon as the sole carbon source Venkatesh Pilli, Ritwik Sarkar PII:
S0272-8842(20)30373-4
DOI:
https://doi.org/10.1016/j.ceramint.2020.02.051
Reference:
CERI 24274
To appear in:
Ceramics International
Received Date: 27 December 2019 Revised Date:
31 January 2020
Accepted Date: 6 February 2020
Please cite this article as: V. Pilli, R. Sarkar, Low carbon containing Al2O3-C refractories with nanocarbon as the sole carbon source, Ceramics International (2020), doi: https://doi.org/10.1016/ j.ceramint.2020.02.051. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier Ltd.
Low carbon containing Al2O3-C refractories with nanocarbon as the sole carbon source Venkatesh Pilli and Ritwik Sarkar Department of Ceramic Engineering, National Institute of Technology, Rourkela – 769008, India E-mail: [email protected] and [email protected]
Abstract: The functional refractories used in the continuous casting process of steel are mostly made up of alumina (Al2O3) - carbon (C) system. Graphite is used as the main carbon source because of the properties it offered. Generally, these refractories contain about 30% residual carbon after coking. But the use of high carbon content leads to several disadvantages like carbon pickup by steel, high heat loss and generation of higher extent of COx gases, etc. Hence, the development of low carbon Al2O3-C refractories without conceding any beneficial properties is the challenge to the refractory technologists. Such a challenge is intended in current study with the use of nanocarbon as a complete alternative for graphite. The variation in physical, mechanical and thermo-mechanical properties with variation in amount of nanocarbon in the composition is studied. Phase analysis and microstructural developments are also evaluated along with the oxidation resistance at different temperatures. Because of its high reactivity, nanocarbon helps to form in-situ ceramic phases at much lower temperatures thereby enhances strength and other properties. But the increase in amount of nanocarbon above 2% found to deteriorate the properties. Keywords: Alumina; graphite; nanocarbon; agglomeration; aluminum carbide; oxidation
Introduction: The flow control devices used in the continuous casting process of steel production, namely: submerged entry nozzles, slide gate plates, ladle shrouds and monoblock stopper rods are generally made-up of alumina-carbon based refractories. Notably, these refractories are important for their thermal, thermos-mechanical and corrosion properties [1–6]. The primary constituent present in this class of refractories is alumina, which is essential to provide mechanical and thermomechanical properties. Addition of carbon provides and enhances properties like thermal shock
resistance owing to its low thermal expansion and superior thermal conductivity, and also provides resistance to corrosion by avoiding the adhesion of slag and metal on the refractory due to its non-wetting nature [7]. In this kind of refractories graphite is used as the principal source of carbon, mainly due to its high oxidation resistance among different carbon sources [8]. The functional refractories of steel casting operations contains residual carbon about 30% after coking [9]. Besides the advantages, existence of carbon in higher amount in the composition associated with several disadvantages also. Particularly at high temperatures, carbon is highly affected to oxidation. Carbon oxidation creates a porosity within the refractory which affects the strength adversely; which can also be penetrated and corroded by slag/steel melts. To control the oxidation of carbon, metal powders like Al and Si are added to the refractory composition. These metal powders acts as antioxidants and form in-situ ceramic phases like Al4C3, SiC, etc. by reacting with carbon at elevated temperatures, that enhances the bonding between refractory aggregates thereby improves strength and thermo-mechanical properties [10–13]. Again, steelmaking is basically a decarburization process. Presence of high carbon content in the refractory linings enhances the chances of carbon pick-up by steel, which affects the properties of steel negatively. It also causes high heat loss due to increased thermal conductivity and also increases the release of high amount of CO and CO2 gases into the atmosphere [8,9,14,15]. In view of all the mentioned lacunas, the carbon content in refractories is to be minimized. To reduce the carbon amount in the composition without scarifying the desired properties; various sources of carbon and their combinations are being studied by the refractory researchers (mostly as partial replacement of graphite) [1,2,16–18]. In the present study, (N220) nanocarbon is used as a sole source of carbon without graphite in alumina-carbon based refractories. The refractories are prepared by conventional alumina-carbon refractory processing techniques. Changes in physical, mechanical and thermomechanical properties is studied by varying nanocarbon content in the composition. X-ray diffraction analysis and microstructural developments and resistance to oxidation of the developed refractories at different amounts of nanocarbon content, at cured and coked at different temperatures are also characterized. The results obtained for the developed compositions are also compared against a reference batch composition containing 25% graphite as carbon source.
Experimental procedure: White tabular alumina of different particle sizes viz. 3-2mm, 1-0.5mm, 0.3-0.05mm and <45µm (Al2O3 - 99.5%, SiO2 ≤ 0.09%, Na2O ≤ 0.40, Fe2O3 ≤ 0.02); Flake graphite of size < 300µm (C 94.1%, volatile matter - 0.8%, ash - 5.08%, sp. surface area - 5.7 m2/gm); N220 nanocarbon, size in the range of 10 - 20nm (C - 98.3%, volatile matter - 1.3%, ash – 0.19%, sp. surface area 120.1 m2/gm); are used as the raw materials. Metal powders of aluminum (Al) and silicon (Si) (size <75 micron, purity >98%); boron carbide (B4C) powder (size <150 micron, purity- 98%) are used as antioxidants. Both powder (fixed carbon ~55.46) and liquid resin (fixed carbon ~39.42); are used as binders. The batch compositions used in the study are detailed in Table 1. Table 1 Batch compositions Prepared Nanocarbon Al N220
Powder resin
Liquid resin
White tabular alumina
Graphite
0
75.0
25.0
-
2.0
1.0
1.0
2.5
4.0
1
99.0
-
1.0
2.0
1.0
1.0
2.5
4.0
2
98.0
-
2.0
2.0
1.0
1.0
2.5
4.0
3
97.0
-
3.0
2.0
1.0
1.0
2.5
4.0
4
96.0
-
4.0
2.0
1.0
1.0
2.5
4.0
Batch
B 4C
Si
The batch compositions are prepared by varying the nanocarbon content from 1 to 4wt% with an increment of 1% without adding graphite. The batch composition made with 25% graphite without nanocarbon is taken as a reference one. The processed batches named after the amount of nanocarbon content present in the composition. Planetary mixer (Hobart N 50) is used to mix the raw materials as per their batch composition. The raw materials are added one after another while mixing; to get a proper homogenous mix. Nanocarbon comprises of nano-sized particles and has very high surface area with obvious agglomeration. The uniform distribution of nanocarbon is essential to get the benefits of its addition. So, firstly nanocarbon is added beside coarser alumina particles, mixed until the nanocarbon coats on the surface of alumina particles, next powder resin, followed by antioxidants, alumina fines, mediums, and graphite. Finally, liquid resin is added and the mixing continued thoroughly till proper consistency is achieved [19]. After mixing, the compositions are pressed into 60mm diameter cylinders and 150mm ×
25mm × 25mm bars using Cr-steel mold in a hydraulic pressing machine (Aimil, India) with a pressure of 150 MPa. Then the pressed shapes are cured at 220oC for 12 hours in a hot air oven (ACMAS, India). Later, the cured samples are coked at different temperatures viz. 1000, 1200, 1400 and 1600oC with a soaking time of 3 hours in an alumina crucible filled with coke particles. Batch preparation is detailed in Figure 1 [19]. Coked and cured samples are tested for several properties related to refractories. Modified Archimedes principle is used to measure bulk density of the samples by vacuum method with kerosene as liquid medium. The cold crushing strength of the cylindrical samples is determined by crushing the samples in a compression-testing machine (Aimil, India). Oxidation resistance of the pre-coked (at 1000oC for 2hr) cylindrical samples of all batch compositions is measured by firing at different temperatures viz. 1000, 1200, 1400 and 1600oC with a soaking time of 3 hours in air atmosphere. The hot modulus rupture of the pre-coked (at 1000oC for 2hr) bar shaped samples of all the compositions is determined by using 3 point bending test at 1400oC in air atmosphere with a soaking time of 30 minutes in a HMOR testing machine (Bysakh & Co., India). Thermal shock resistance of the pre-coked (at 1000oC for 2hr) cylindrical samples of all the compositions is evaluated by heating the samples in coke embedded condition to 1100oC at a heating rate of 5oC/min with a holding time of 30 minutes, followed by quenching them into water at room temperature. The same thermal cycles is repeated for three times. The compressive strength of the samples is measured after each thermal shock cycle. The percentage of residual strength retained after thermal shocks against the original strength of the samples is represented as its thermal shock resistance. The phase analysis of the coked samples is studied by using Cu Kα radiation through Ni filter, in an X-ray diffractometer (Rigaku, Japan). Microstructural developments in coked samples (fractured surface) are investigated by using a field emission scanning electron microscope (Nova nanoSEM 450, FEI, US) with the support of energy dispersive X-ray analysis (SDD EDS detector, Bruker, Germany).
Figure 1 Process flowsheet
Results and discussions: Phase analysis: The phase analysis is carried out in the matrix portion only (to avoid the interference of the coarser alumina aggregates) of all the processed batch compositions coked at different temperatures. The X-Ray diffraction patterns are provided in Figures 2-6. Phase analysis data of 25% graphite comprising batch is provided in Figure 2, showing graphite as the major phase along with the peaks of alumina at all coking temperatures due to the existence of high amount of graphite flakes in the matrix part of the composition. No other bonding phases are found. The higher amount of graphite flakes distributed well in the composition might restrict the bonding among refractory aggregates present in the system. However, in the case of nanocarbon batches, provided in Figures 3-6, alumina as the major phase along with smaller peaks of aluminum carbide and aluminum borate at temperatures 1000 and 1200oC. Increase in coking temperature to 1400 and 1600oC resulted in the disappearance of the carbide and borate phases, and the smaller peaks of silicon carbide have appeared.
Figure 2 XRD analysis of 25% graphite batch
Figure 3 XRD analysis of 1% nanocarbon batch
Figure 4 XRD analysis of 2% nanocarbon batch
Figure 5 XRD analysis diffraction of 3% nanocarbon batch
Figure 6 XRD analysis of 4% nanocarbon batch
Nanocarbon possesses high reactivity compared to the graphite flake because of its smaller particle size and high surface area, can easily form the ceramic bonding phases. The formation of aluminum borate phase in the nanocarbon containing batch compositions is due to the addition of boron carbide. It reacts with CO or O2 present in the system and forms B2O3 liquid at 450oC, according to Equations 1 and 2; resulting in the formation of a thin transparent film that helps to inhibit oxidation. Further, increase in temperature to 1000 and 1200oC it reacts with fine alumina particle present in the system and forms aluminum borate according to Equation 3 [1,20–24]. + 6
( ) ( )
+ 4
2
( )
( ) ( )
= 2
( )
= 2
+ 2
( )
( )
+7
( ) −
+
−−−−−
( ) −
=
( )
.1
−−−−−
−−−−−−
.2 . 3
The Al metal powder starts melting about 660oC; when the temperature is increased to 1000oC, Al reacts with surrounding carbon and forms aluminum carbide according to Equations 4 and 5. 4 8
(, )
(, )
+3
+6
( )
( )
=
= 2
( ) ( )
−−−−−−
+3
( )
.4
−−−−−−
.5
But, when the temperature is increased to 1400 and 1600oC, the carbide phase disappears, as it completely disintegrated into alumina and carbon according to the Equation 6. This also adversely affects the mechanical properties of the compositions. ( )
+6
( )
= 2
( )
+9
( )
−−−−−−
.6
The formation of silicon carbide started at 1400oC due to the presence of Si metal powder. The nanocarbon particles react with SiO vapor and form silicon carbide according to the Equations 7 and 8. 2#$( ) + #$
( )
+2
( )
( )
= 2#$
= #$
( )
( )
+
−−−−−− ( )
−−−−−−
.7
.8
The formation of in-situ ceramic bonding phases in nanocarbon containing compositions helps to improve the mechanical properties [1–3,20,21,25], as observed during the CCS study.
Microstructural analysis: Microstructural study is carried out on the fractured surface of the samples. The micrographs of the samples of all the compositions coked at different temperatures are provided in Figures 7-10. The micrographs of 25% graphite containing composition coked at 1000oC is provided in Figure 7 along with the EDX analysis showing sheet like structures, which is graphite [1]. The EDX spectrum also displays the presence of carbon with small amount of oxygen and aluminum.
Figure 7 25% graphite containing composition coked at 1000oC All the other nanocarbon containing compositions showing column or needle-like structure at temperatures 1000 and 1200oC is due to the formation of aluminum carbide [12,26,27] provided in Figure 8. No such carbide formations are observed in 25% graphite containing batch composition. The graphite flakes are much larger in size compared to nanocarbon particles, less reactive and so the in-situ ceramic phases like carbides formation at lower temperatures is negligible or very limited. Nanocarbon particles are extremely reactive due to their fineness, can initiate the reaction for carbide phase formation at much lower temperatures, as shown in Figure 8, compared to the graphite flakes. The micrographs of the nanocarbon containing compositions coked at 1600oC displayed more compact microstructure than that of the traditional graphite containing one, provided in Figure 9. The graphite containing composition shows flakes of graphite and few alumina aggregates. The 1% nanocarbon batch coked at 1600oC showed more compact microstructure than the graphite containing and 4% nanocarbon-containing batches. Much finer nanocarbon particles can effectively fill the inter particle voids present in the refractory comparative to coarser graphite
particles and thus improves the packing and compactness. But, at higher amount of nanocarbon, presence of excess amount of nanocarbon particles gets agglomerated due to their surface charges, does not fill the voids properly and remains free in the system, resulting in the generation of porosity, which affects the strength of the refractory [1,28].
a
c
b
7
cps/eV B-K C-K
O-K
Al-K
Si-KA
O
Al
Si
Au-MAB
6
5
4
3
B
Au C
Au
2
1
0 0.5
1.0
1.5
2.0
2.5
3.0
keV
Figure 8 Micrographs of a) 1NC b) 3NC coked at 1000oC c) 2NC coked at 1200oC along with EDX spectrum
a
b
c
Figure 9 Micrographs of a) 25G b) 1NC c) 4NC compositions coked at 1600oC
The micrograph of 4% nanocarbon composition coked at 1000oC provided in Figure 10, showing formation of aluminum carbide on the surface of alumina grain. The uniform dispersion of nanocarbon in the composition is important to achieve the desired properties [19]. To distribute nanocarbon in the composition, firstly the nanocarbon is mixed with the coarse alumina particles and then other materials are added one after the other as detailed in Figure 1. The nanocarbon particles present as a coat on the alumina grain surface react with aluminum metal powder or finer alumina particles also present in the matrix phase, as a coating on the alumina grain and form aluminum carbide. Thus, the bonding within the composition and in between alumina particles also increases through this carbide formation.
Figure 10 Formation of aluminum carbide on alumina grain in 4% NC composition coked at 1000oC
Bulk density: The bulk density values of all the batch compositions cured and coked at different temperatures are provided in Figure 11. The bulk density of 25% graphite batch is low compared to all the other compositions due to the presence of high amount graphite (carbon source), low-density material (relatively). For the other batch compositions, graphite is mostly replaced by alumina
Figure 11 Bulk density of the compositions
(having much higher specific gravity than graphite) and nanocarbon, showing comparatively much higher density values. In nanocarbon containing compositions, the density values are decreasing marginally and gradually with the increase in amount of nanocarbon in the batch composition. Increase in the amount of nanocarbon decreases the alumina content; replacing a higher density material (alumina) with lower density material (carbon) causes reduction in density. Also, for compositions with higher amount of nanocarbon, all the carbon particles do not enter into the inter particle voids of other refractory aggregates and remains free in the system which may get oxidized during coking, resulting in reduction in density [1,10,28]. Also, the bulk density values for all the compositions are found to decrease at higher coking temperatures. In-situ ceramic phase formation, especially aluminum carbide, dissociates at high temperature, resulting in the loss in extra bonding phase that developed on coking at 1000 and 1200oC, thus the bonding weakens and results in reduction in density values.
Cold crushing strength: The cold crushing strength values of the processed batch compositions, cured and coked at different temperatures, are provided in Figure 12. Nanocarbon containing compositions show higher strength values for all the temperatures of processing. Fine nanocarbon particles fill the voids in the refractory aggregate system better than the flaky and coarse graphite particles, thus improves the compactness and strength for all temperatures. Also, nanocarbon particles are reactive in nature due to their extreme fineness, thus they form in-situ ceramic phases easily, as found in the phase analysis and microstructural study, which provide extra bonding in the composition and enhances the strength further. For higher amount of nanocarbon in the composition, as the filling of voids is affected due to agglomeration, the strength values marginally decreases with increasing amount of nanocarbon content. Again, this carbide formation is negligible or very limited in case of graphite containing composition, due to much coarser less reactive graphite flake particles, so such enhancement of strength is also not prominent in the composition. At higher coking temperatures, disintegration of aluminum carbide occurs forming alumina and carbon [1,2,20,21], as also observed in phase analysis study. Thus the extra bonding, especially for the nanocarbon containing compositions, is lost at higher coking temperatures and results in gradual reduction in strength. But as the carbide formation is
not that effective in graphite containing reference batch, this degradation in strength is not observed / prominent in the composition.
Figure 12 Cold crushing strength of the compositions
Hot modulus of rupture (HMOR): Hot modulus of rupture (HMOR) test is performed on the bar shaped samples coked at 1000oC of all the compositions by 3-point bending technique at 1400oC in air atmosphere. The HMOR strength values are provided in Figure 13. The HMOR value increases from 1% to 2% nanocarbon containing batch composition but further increase in the amount of nanocarbon does not improve the hot strength. Because of its high surface area, nanocarbon possesses high reactivity, can form greater extent of in-situ ceramic phases in the matrix phase of the compositions. The formation of in-situ ceramic phases like aluminum carbide and small amount of silicon carbide is responsible for providing extra bonding in the compositions at higher temperatures and the improvement in the strength values. As the test is performed in the oxidizing atmosphere, the loss of carbon through oxidation affects the strength. The high amount of nanocarbon particles remain free in the system, which are prone to oxidize faster compared to the graphite flake because of its high surface area. As a result, higher nanocarbon containing compositions showed degradation in hot strength. In 25%
graphite containing batch the formation of in-situ phases is negligible due to coarser size of graphite particles and the presence of higher amount of graphite increases the tendency of oxidation at higher temperatures. Hence, results in a lower hot strength values.
Figure 13 Hot modulus of rupture of the compositions
Oxidation resistance: Oxidation resistance is measured by firing the samples at different temperatures viz. 1000, 1200, 1400 and 1600oC of all the batch compositions in air atmosphere with a soaking time of 3 hours. The samples are pre-coked at 1000oC for 2 hours before oxidation. The loss of carbon in the samples after oxidation resulted in weight loss and porosity increment. Hence the carbon loss occurred in the samples is related to the weight loss. The percentage of carbon loss in the compositions after oxidation is provided in Figure 14. The carbon loss increases with increase in temperature in all the cases. The total carbon content in the 25% graphite-containing batch is about 28% (Graphite + Powder resin + Liquid resin) and for the nanocarbon batches. For, 1NC to 4NC batches the carbon content varies between 4 to 7% respectively. The percent carbon loss is lower in the case of 25% graphite containing batch composition at 1000 and 1200oC compared to other batch compositions processed. The nanocarbon having smaller particle size and high surface area is prone to oxidize faster than the graphite flake. Further increase in firing temperature, the rate of oxidation in 25% graphite containing batch increased compared to 1NC and 2NC batches. The rate of oxidation in other nanocarbon batches
viz. 3NC and 4NC is more due to the excess amount of nanocarbon present in it. Higher amount of nanocarbon particles does not enter into the gaps /pores between other refractory aggregates and remains free in the system, which can be oxidized easily. The photographs of samples of all the batch compositions (cut surface) after oxidation are provided in Figures 15-19.
Figure 14 Carbon loss from the compositions after oxidation
Figure 15 25% Graphite batch after oxidation at 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Figure 16 1% Nanocarbon batch after oxidation at 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Figure 17 2% Nanocarbon batch after oxidation 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Figure 18 3% Nanocarbon batch after oxidation at 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Figure 19 4% Nanocarbon batch after oxidation at 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Outer decarburized layer is depleted with carbon and contains only alumina. This only alumina part densify at high temperature during firing The bulk density of the outer layer is measured for all the batch compositions after oxidation at 1400 and 1600oC, which is provided in Figure 20. Decarburized layer of the 25% graphite-containing batch does not densify properly and retain any strength after the oxidation due to the formation of much larger sized (compared to nanocarbon) pores from graphite particles in higher amount. These larger pores do not annihilated and do not allow the layer to densify. In 25% graphite batch, higher amount of graphite particles distribute well among the batch composition restricts the direct bonding between alumina particles; and leaves large pores after oxidation. However, the nanocarboncontaining compositions showed higher density values in the decarburized layer even after the oxidation due to the sintering of alumina particles. Nano-sized pores generated on oxidation of nano-sized particles of nanocarbon get removed from the material at higher temperature and cause greater densification and also corresponding strength in the decarburized layer. This densified decarburized layer does not allow oxygen to enter inside the refractory composition and thus prevents further oxidation of carbon particles in the inner layers. But with increasing amount of nanocarbon content in the composition, as the amount of finer pores are increasing in the decarburized layer, the density values fall gradually and marginally.
Figure 20 Bulk density of the oxidized layer
Thermal shock resistance: The thermal shock resistance for all the batch compositions, pre-coked at 1000oC for 2 hours, is evaluated by heating them at 1100oC in coked condition and water quenching method. The percentage of retained strength values after thermal shock of all the compositions is provided in Figure 21. It is observed that the retained strength values of all the compositions are decreasing with the increase in the number of thermal shock cycles.
Figure 21 Retained strength after thermal shock The sudden change in temperature due to thermal shock leads to the generation of thermal stresses within the samples, leading to cracks that deteriorates the strength. The 1% nanocarboncontaining batch showed better retention of strength compared to all other batches processed. Even the 2% nanocarbon composition batch showed almost an equal value with the reference batch composition consists of 25% graphite. These nanocarbon particles can absorb and relieve the stresses generated during thermal expansion and shrinkage of the refractory thereby reduces the misdistribution stress within the sample and results in the improvement of thermal shock resistance [15, 29]. Tamura et al. reported [29] that 1.5% nanocarbon batch showed equivalent thermal spalling properties to 18% graphite containing composition in MgO-C system. The batch
compositions with more than 2% nanocarbon showed less residual strength might be due to greater extent of oxidation of the agglomerated nanocarbon particles.
Conclusion: The batch compositions prepared with nanocarbon showed higher bulk density and cold crushing strength compared to the reference 25% graphite containing composition. The graphite content of the reference batch is replaced by fine alumina and nanocarbon in nanocarbon containing batches. Replacing lower density material (graphite) with higher density material (alumina) enhances density values. Because of its smaller size, nanocarbon can easily fill the gaps among the different sized aggregate fractions of the refractory thereby improves density. The strength values of the nanocarbon batches increases up to a coking temperature of 1200oC due to the formation of aluminum carbide phase but decreases at higher coking temperatures due to its dissociation. Nanocarbon is active due to its high surface area can initiate the reaction at much lower temperatures than the graphite flake. Oxidation resistance is measured in terms of weight loss, and the corresponding carbon loss is measured. The 25% graphite batch after decarburization due to oxidation does not retain any strength. Graphite flake has larger size compared to nanocarbon, which on oxidation leaves a bigger pore, does not allow the remaining alumina to sinter and resulting in weak structure. However, the nanocarbon batches show higher density in decarburized layer after oxidation due to the sintering of alumina particles in the composition. Dense decarburized layer will restrict further oxidation. The residual strength values are observed to decrease in all the compositions with the increase in number of thermal spalling cycles. The 1% nanocarbon-containing batch showed better retention of strength compared to all the other compositions. The increase in amount of nanocarbon content above 2% does not promote the properties any further. Acknowledgement: The authors thankfully acknowledge the financial support of CSIR, GOI for sponsored project, Grant-in-Aid, Ref No. 22(0767)/18/EMR-II. The authors also acknowledge the support of the staff of Department of Ceramic Engg, NIT, Rourkela for their help in experimentation; Mr. S. K. Shrivastava and Mr. B.C. Nanda of IFGL Refractories, India, for their support in sample preparation. Also special thanks to Mr. S. Chatterjee, Almatis, India for support in HMOR testing.
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Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:
S0272-8842(20)30373-4
DOI:
https://doi.org/10.1016/j.ceramint.2020.02.051
Reference:
CERI 24274
To appear in:
Ceramics International
Received Date: 27 December 2019 Revised Date:
31 January 2020
Accepted Date: 6 February 2020
Please cite this article as: V. Pilli, R. Sarkar, Low carbon containing Al2O3-C refractories with nanocarbon as the sole carbon source, Ceramics International (2020), doi: https://doi.org/10.1016/ j.ceramint.2020.02.051. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier Ltd.
Low carbon containing Al2O3-C refractories with nanocarbon as the sole carbon source Venkatesh Pilli and Ritwik Sarkar Department of Ceramic Engineering, National Institute of Technology, Rourkela – 769008, India E-mail: [email protected] and [email protected]
Abstract: The functional refractories used in the continuous casting process of steel are mostly made up of alumina (Al2O3) - carbon (C) system. Graphite is used as the main carbon source because of the properties it offered. Generally, these refractories contain about 30% residual carbon after coking. But the use of high carbon content leads to several disadvantages like carbon pickup by steel, high heat loss and generation of higher extent of COx gases, etc. Hence, the development of low carbon Al2O3-C refractories without conceding any beneficial properties is the challenge to the refractory technologists. Such a challenge is intended in current study with the use of nanocarbon as a complete alternative for graphite. The variation in physical, mechanical and thermo-mechanical properties with variation in amount of nanocarbon in the composition is studied. Phase analysis and microstructural developments are also evaluated along with the oxidation resistance at different temperatures. Because of its high reactivity, nanocarbon helps to form in-situ ceramic phases at much lower temperatures thereby enhances strength and other properties. But the increase in amount of nanocarbon above 2% found to deteriorate the properties. Keywords: Alumina; graphite; nanocarbon; agglomeration; aluminum carbide; oxidation
Introduction: The flow control devices used in the continuous casting process of steel production, namely: submerged entry nozzles, slide gate plates, ladle shrouds and monoblock stopper rods are generally made-up of alumina-carbon based refractories. Notably, these refractories are important for their thermal, thermos-mechanical and corrosion properties [1–6]. The primary constituent present in this class of refractories is alumina, which is essential to provide mechanical and thermomechanical properties. Addition of carbon provides and enhances properties like thermal shock
resistance owing to its low thermal expansion and superior thermal conductivity, and also provides resistance to corrosion by avoiding the adhesion of slag and metal on the refractory due to its non-wetting nature [7]. In this kind of refractories graphite is used as the principal source of carbon, mainly due to its high oxidation resistance among different carbon sources [8]. The functional refractories of steel casting operations contains residual carbon about 30% after coking [9]. Besides the advantages, existence of carbon in higher amount in the composition associated with several disadvantages also. Particularly at high temperatures, carbon is highly affected to oxidation. Carbon oxidation creates a porosity within the refractory which affects the strength adversely; which can also be penetrated and corroded by slag/steel melts. To control the oxidation of carbon, metal powders like Al and Si are added to the refractory composition. These metal powders acts as antioxidants and form in-situ ceramic phases like Al4C3, SiC, etc. by reacting with carbon at elevated temperatures, that enhances the bonding between refractory aggregates thereby improves strength and thermo-mechanical properties [10–13]. Again, steelmaking is basically a decarburization process. Presence of high carbon content in the refractory linings enhances the chances of carbon pick-up by steel, which affects the properties of steel negatively. It also causes high heat loss due to increased thermal conductivity and also increases the release of high amount of CO and CO2 gases into the atmosphere [8,9,14,15]. In view of all the mentioned lacunas, the carbon content in refractories is to be minimized. To reduce the carbon amount in the composition without scarifying the desired properties; various sources of carbon and their combinations are being studied by the refractory researchers (mostly as partial replacement of graphite) [1,2,16–18]. In the present study, (N220) nanocarbon is used as a sole source of carbon without graphite in alumina-carbon based refractories. The refractories are prepared by conventional alumina-carbon refractory processing techniques. Changes in physical, mechanical and thermomechanical properties is studied by varying nanocarbon content in the composition. X-ray diffraction analysis and microstructural developments and resistance to oxidation of the developed refractories at different amounts of nanocarbon content, at cured and coked at different temperatures are also characterized. The results obtained for the developed compositions are also compared against a reference batch composition containing 25% graphite as carbon source.
Experimental procedure: White tabular alumina of different particle sizes viz. 3-2mm, 1-0.5mm, 0.3-0.05mm and <45µm (Al2O3 - 99.5%, SiO2 ≤ 0.09%, Na2O ≤ 0.40, Fe2O3 ≤ 0.02); Flake graphite of size < 300µm (C 94.1%, volatile matter - 0.8%, ash - 5.08%, sp. surface area - 5.7 m2/gm); N220 nanocarbon, size in the range of 10 - 20nm (C - 98.3%, volatile matter - 1.3%, ash – 0.19%, sp. surface area 120.1 m2/gm); are used as the raw materials. Metal powders of aluminum (Al) and silicon (Si) (size <75 micron, purity >98%); boron carbide (B4C) powder (size <150 micron, purity- 98%) are used as antioxidants. Both powder (fixed carbon ~55.46) and liquid resin (fixed carbon ~39.42); are used as binders. The batch compositions used in the study are detailed in Table 1. Table 1 Batch compositions Prepared Nanocarbon Al N220
Powder resin
Liquid resin
White tabular alumina
Graphite
0
75.0
25.0
-
2.0
1.0
1.0
2.5
4.0
1
99.0
-
1.0
2.0
1.0
1.0
2.5
4.0
2
98.0
-
2.0
2.0
1.0
1.0
2.5
4.0
3
97.0
-
3.0
2.0
1.0
1.0
2.5
4.0
4
96.0
-
4.0
2.0
1.0
1.0
2.5
4.0
Batch
B 4C
Si
The batch compositions are prepared by varying the nanocarbon content from 1 to 4wt% with an increment of 1% without adding graphite. The batch composition made with 25% graphite without nanocarbon is taken as a reference one. The processed batches named after the amount of nanocarbon content present in the composition. Planetary mixer (Hobart N 50) is used to mix the raw materials as per their batch composition. The raw materials are added one after another while mixing; to get a proper homogenous mix. Nanocarbon comprises of nano-sized particles and has very high surface area with obvious agglomeration. The uniform distribution of nanocarbon is essential to get the benefits of its addition. So, firstly nanocarbon is added beside coarser alumina particles, mixed until the nanocarbon coats on the surface of alumina particles, next powder resin, followed by antioxidants, alumina fines, mediums, and graphite. Finally, liquid resin is added and the mixing continued thoroughly till proper consistency is achieved [19]. After mixing, the compositions are pressed into 60mm diameter cylinders and 150mm ×
25mm × 25mm bars using Cr-steel mold in a hydraulic pressing machine (Aimil, India) with a pressure of 150 MPa. Then the pressed shapes are cured at 220oC for 12 hours in a hot air oven (ACMAS, India). Later, the cured samples are coked at different temperatures viz. 1000, 1200, 1400 and 1600oC with a soaking time of 3 hours in an alumina crucible filled with coke particles. Batch preparation is detailed in Figure 1 [19]. Coked and cured samples are tested for several properties related to refractories. Modified Archimedes principle is used to measure bulk density of the samples by vacuum method with kerosene as liquid medium. The cold crushing strength of the cylindrical samples is determined by crushing the samples in a compression-testing machine (Aimil, India). Oxidation resistance of the pre-coked (at 1000oC for 2hr) cylindrical samples of all batch compositions is measured by firing at different temperatures viz. 1000, 1200, 1400 and 1600oC with a soaking time of 3 hours in air atmosphere. The hot modulus rupture of the pre-coked (at 1000oC for 2hr) bar shaped samples of all the compositions is determined by using 3 point bending test at 1400oC in air atmosphere with a soaking time of 30 minutes in a HMOR testing machine (Bysakh & Co., India). Thermal shock resistance of the pre-coked (at 1000oC for 2hr) cylindrical samples of all the compositions is evaluated by heating the samples in coke embedded condition to 1100oC at a heating rate of 5oC/min with a holding time of 30 minutes, followed by quenching them into water at room temperature. The same thermal cycles is repeated for three times. The compressive strength of the samples is measured after each thermal shock cycle. The percentage of residual strength retained after thermal shocks against the original strength of the samples is represented as its thermal shock resistance. The phase analysis of the coked samples is studied by using Cu Kα radiation through Ni filter, in an X-ray diffractometer (Rigaku, Japan). Microstructural developments in coked samples (fractured surface) are investigated by using a field emission scanning electron microscope (Nova nanoSEM 450, FEI, US) with the support of energy dispersive X-ray analysis (SDD EDS detector, Bruker, Germany).
Figure 1 Process flowsheet
Results and discussions: Phase analysis: The phase analysis is carried out in the matrix portion only (to avoid the interference of the coarser alumina aggregates) of all the processed batch compositions coked at different temperatures. The X-Ray diffraction patterns are provided in Figures 2-6. Phase analysis data of 25% graphite comprising batch is provided in Figure 2, showing graphite as the major phase along with the peaks of alumina at all coking temperatures due to the existence of high amount of graphite flakes in the matrix part of the composition. No other bonding phases are found. The higher amount of graphite flakes distributed well in the composition might restrict the bonding among refractory aggregates present in the system. However, in the case of nanocarbon batches, provided in Figures 3-6, alumina as the major phase along with smaller peaks of aluminum carbide and aluminum borate at temperatures 1000 and 1200oC. Increase in coking temperature to 1400 and 1600oC resulted in the disappearance of the carbide and borate phases, and the smaller peaks of silicon carbide have appeared.
Figure 2 XRD analysis of 25% graphite batch
Figure 3 XRD analysis of 1% nanocarbon batch
Figure 4 XRD analysis of 2% nanocarbon batch
Figure 5 XRD analysis diffraction of 3% nanocarbon batch
Figure 6 XRD analysis of 4% nanocarbon batch
Nanocarbon possesses high reactivity compared to the graphite flake because of its smaller particle size and high surface area, can easily form the ceramic bonding phases. The formation of aluminum borate phase in the nanocarbon containing batch compositions is due to the addition of boron carbide. It reacts with CO or O2 present in the system and forms B2O3 liquid at 450oC, according to Equations 1 and 2; resulting in the formation of a thin transparent film that helps to inhibit oxidation. Further, increase in temperature to 1000 and 1200oC it reacts with fine alumina particle present in the system and forms aluminum borate according to Equation 3 [1,20–24]. + 6
( ) ( )
+ 4
2
( )
( ) ( )
= 2
( )
= 2
+ 2
( )
( )
+7
( ) −
+
−−−−−
( ) −
=
( )
.1
−−−−−
−−−−−−
.2 . 3
The Al metal powder starts melting about 660oC; when the temperature is increased to 1000oC, Al reacts with surrounding carbon and forms aluminum carbide according to Equations 4 and 5. 4 8
(, )
(, )
+3
+6
( )
( )
=
= 2
( ) ( )
−−−−−−
+3
( )
.4
−−−−−−
.5
But, when the temperature is increased to 1400 and 1600oC, the carbide phase disappears, as it completely disintegrated into alumina and carbon according to the Equation 6. This also adversely affects the mechanical properties of the compositions. ( )
+6
( )
= 2
( )
+9
( )
−−−−−−
.6
The formation of silicon carbide started at 1400oC due to the presence of Si metal powder. The nanocarbon particles react with SiO vapor and form silicon carbide according to the Equations 7 and 8. 2#$( ) + #$
( )
+2
( )
( )
= 2#$
= #$
( )
( )
+
−−−−−− ( )
−−−−−−
.7
.8
The formation of in-situ ceramic bonding phases in nanocarbon containing compositions helps to improve the mechanical properties [1–3,20,21,25], as observed during the CCS study.
Microstructural analysis: Microstructural study is carried out on the fractured surface of the samples. The micrographs of the samples of all the compositions coked at different temperatures are provided in Figures 7-10. The micrographs of 25% graphite containing composition coked at 1000oC is provided in Figure 7 along with the EDX analysis showing sheet like structures, which is graphite [1]. The EDX spectrum also displays the presence of carbon with small amount of oxygen and aluminum.
Figure 7 25% graphite containing composition coked at 1000oC All the other nanocarbon containing compositions showing column or needle-like structure at temperatures 1000 and 1200oC is due to the formation of aluminum carbide [12,26,27] provided in Figure 8. No such carbide formations are observed in 25% graphite containing batch composition. The graphite flakes are much larger in size compared to nanocarbon particles, less reactive and so the in-situ ceramic phases like carbides formation at lower temperatures is negligible or very limited. Nanocarbon particles are extremely reactive due to their fineness, can initiate the reaction for carbide phase formation at much lower temperatures, as shown in Figure 8, compared to the graphite flakes. The micrographs of the nanocarbon containing compositions coked at 1600oC displayed more compact microstructure than that of the traditional graphite containing one, provided in Figure 9. The graphite containing composition shows flakes of graphite and few alumina aggregates. The 1% nanocarbon batch coked at 1600oC showed more compact microstructure than the graphite containing and 4% nanocarbon-containing batches. Much finer nanocarbon particles can effectively fill the inter particle voids present in the refractory comparative to coarser graphite
particles and thus improves the packing and compactness. But, at higher amount of nanocarbon, presence of excess amount of nanocarbon particles gets agglomerated due to their surface charges, does not fill the voids properly and remains free in the system, resulting in the generation of porosity, which affects the strength of the refractory [1,28].
a
c
b
7
cps/eV B-K C-K
O-K
Al-K
Si-KA
O
Al
Si
Au-MAB
6
5
4
3
B
Au C
Au
2
1
0 0.5
1.0
1.5
2.0
2.5
3.0
keV
Figure 8 Micrographs of a) 1NC b) 3NC coked at 1000oC c) 2NC coked at 1200oC along with EDX spectrum
a
b
c
Figure 9 Micrographs of a) 25G b) 1NC c) 4NC compositions coked at 1600oC
The micrograph of 4% nanocarbon composition coked at 1000oC provided in Figure 10, showing formation of aluminum carbide on the surface of alumina grain. The uniform dispersion of nanocarbon in the composition is important to achieve the desired properties [19]. To distribute nanocarbon in the composition, firstly the nanocarbon is mixed with the coarse alumina particles and then other materials are added one after the other as detailed in Figure 1. The nanocarbon particles present as a coat on the alumina grain surface react with aluminum metal powder or finer alumina particles also present in the matrix phase, as a coating on the alumina grain and form aluminum carbide. Thus, the bonding within the composition and in between alumina particles also increases through this carbide formation.
Figure 10 Formation of aluminum carbide on alumina grain in 4% NC composition coked at 1000oC
Bulk density: The bulk density values of all the batch compositions cured and coked at different temperatures are provided in Figure 11. The bulk density of 25% graphite batch is low compared to all the other compositions due to the presence of high amount graphite (carbon source), low-density material (relatively). For the other batch compositions, graphite is mostly replaced by alumina
Figure 11 Bulk density of the compositions
(having much higher specific gravity than graphite) and nanocarbon, showing comparatively much higher density values. In nanocarbon containing compositions, the density values are decreasing marginally and gradually with the increase in amount of nanocarbon in the batch composition. Increase in the amount of nanocarbon decreases the alumina content; replacing a higher density material (alumina) with lower density material (carbon) causes reduction in density. Also, for compositions with higher amount of nanocarbon, all the carbon particles do not enter into the inter particle voids of other refractory aggregates and remains free in the system which may get oxidized during coking, resulting in reduction in density [1,10,28]. Also, the bulk density values for all the compositions are found to decrease at higher coking temperatures. In-situ ceramic phase formation, especially aluminum carbide, dissociates at high temperature, resulting in the loss in extra bonding phase that developed on coking at 1000 and 1200oC, thus the bonding weakens and results in reduction in density values.
Cold crushing strength: The cold crushing strength values of the processed batch compositions, cured and coked at different temperatures, are provided in Figure 12. Nanocarbon containing compositions show higher strength values for all the temperatures of processing. Fine nanocarbon particles fill the voids in the refractory aggregate system better than the flaky and coarse graphite particles, thus improves the compactness and strength for all temperatures. Also, nanocarbon particles are reactive in nature due to their extreme fineness, thus they form in-situ ceramic phases easily, as found in the phase analysis and microstructural study, which provide extra bonding in the composition and enhances the strength further. For higher amount of nanocarbon in the composition, as the filling of voids is affected due to agglomeration, the strength values marginally decreases with increasing amount of nanocarbon content. Again, this carbide formation is negligible or very limited in case of graphite containing composition, due to much coarser less reactive graphite flake particles, so such enhancement of strength is also not prominent in the composition. At higher coking temperatures, disintegration of aluminum carbide occurs forming alumina and carbon [1,2,20,21], as also observed in phase analysis study. Thus the extra bonding, especially for the nanocarbon containing compositions, is lost at higher coking temperatures and results in gradual reduction in strength. But as the carbide formation is
not that effective in graphite containing reference batch, this degradation in strength is not observed / prominent in the composition.
Figure 12 Cold crushing strength of the compositions
Hot modulus of rupture (HMOR): Hot modulus of rupture (HMOR) test is performed on the bar shaped samples coked at 1000oC of all the compositions by 3-point bending technique at 1400oC in air atmosphere. The HMOR strength values are provided in Figure 13. The HMOR value increases from 1% to 2% nanocarbon containing batch composition but further increase in the amount of nanocarbon does not improve the hot strength. Because of its high surface area, nanocarbon possesses high reactivity, can form greater extent of in-situ ceramic phases in the matrix phase of the compositions. The formation of in-situ ceramic phases like aluminum carbide and small amount of silicon carbide is responsible for providing extra bonding in the compositions at higher temperatures and the improvement in the strength values. As the test is performed in the oxidizing atmosphere, the loss of carbon through oxidation affects the strength. The high amount of nanocarbon particles remain free in the system, which are prone to oxidize faster compared to the graphite flake because of its high surface area. As a result, higher nanocarbon containing compositions showed degradation in hot strength. In 25%
graphite containing batch the formation of in-situ phases is negligible due to coarser size of graphite particles and the presence of higher amount of graphite increases the tendency of oxidation at higher temperatures. Hence, results in a lower hot strength values.
Figure 13 Hot modulus of rupture of the compositions
Oxidation resistance: Oxidation resistance is measured by firing the samples at different temperatures viz. 1000, 1200, 1400 and 1600oC of all the batch compositions in air atmosphere with a soaking time of 3 hours. The samples are pre-coked at 1000oC for 2 hours before oxidation. The loss of carbon in the samples after oxidation resulted in weight loss and porosity increment. Hence the carbon loss occurred in the samples is related to the weight loss. The percentage of carbon loss in the compositions after oxidation is provided in Figure 14. The carbon loss increases with increase in temperature in all the cases. The total carbon content in the 25% graphite-containing batch is about 28% (Graphite + Powder resin + Liquid resin) and for the nanocarbon batches. For, 1NC to 4NC batches the carbon content varies between 4 to 7% respectively. The percent carbon loss is lower in the case of 25% graphite containing batch composition at 1000 and 1200oC compared to other batch compositions processed. The nanocarbon having smaller particle size and high surface area is prone to oxidize faster than the graphite flake. Further increase in firing temperature, the rate of oxidation in 25% graphite containing batch increased compared to 1NC and 2NC batches. The rate of oxidation in other nanocarbon batches
viz. 3NC and 4NC is more due to the excess amount of nanocarbon present in it. Higher amount of nanocarbon particles does not enter into the gaps /pores between other refractory aggregates and remains free in the system, which can be oxidized easily. The photographs of samples of all the batch compositions (cut surface) after oxidation are provided in Figures 15-19.
Figure 14 Carbon loss from the compositions after oxidation
Figure 15 25% Graphite batch after oxidation at 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Figure 16 1% Nanocarbon batch after oxidation at 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Figure 17 2% Nanocarbon batch after oxidation 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Figure 18 3% Nanocarbon batch after oxidation at 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Figure 19 4% Nanocarbon batch after oxidation at 1) 1000oC, 2) 1200oC, 3) 1400oC and 4) 1600oC
Outer decarburized layer is depleted with carbon and contains only alumina. This only alumina part densify at high temperature during firing The bulk density of the outer layer is measured for all the batch compositions after oxidation at 1400 and 1600oC, which is provided in Figure 20. Decarburized layer of the 25% graphite-containing batch does not densify properly and retain any strength after the oxidation due to the formation of much larger sized (compared to nanocarbon) pores from graphite particles in higher amount. These larger pores do not annihilated and do not allow the layer to densify. In 25% graphite batch, higher amount of graphite particles distribute well among the batch composition restricts the direct bonding between alumina particles; and leaves large pores after oxidation. However, the nanocarboncontaining compositions showed higher density values in the decarburized layer even after the oxidation due to the sintering of alumina particles. Nano-sized pores generated on oxidation of nano-sized particles of nanocarbon get removed from the material at higher temperature and cause greater densification and also corresponding strength in the decarburized layer. This densified decarburized layer does not allow oxygen to enter inside the refractory composition and thus prevents further oxidation of carbon particles in the inner layers. But with increasing amount of nanocarbon content in the composition, as the amount of finer pores are increasing in the decarburized layer, the density values fall gradually and marginally.
Figure 20 Bulk density of the oxidized layer
Thermal shock resistance: The thermal shock resistance for all the batch compositions, pre-coked at 1000oC for 2 hours, is evaluated by heating them at 1100oC in coked condition and water quenching method. The percentage of retained strength values after thermal shock of all the compositions is provided in Figure 21. It is observed that the retained strength values of all the compositions are decreasing with the increase in the number of thermal shock cycles.
Figure 21 Retained strength after thermal shock The sudden change in temperature due to thermal shock leads to the generation of thermal stresses within the samples, leading to cracks that deteriorates the strength. The 1% nanocarboncontaining batch showed better retention of strength compared to all other batches processed. Even the 2% nanocarbon composition batch showed almost an equal value with the reference batch composition consists of 25% graphite. These nanocarbon particles can absorb and relieve the stresses generated during thermal expansion and shrinkage of the refractory thereby reduces the misdistribution stress within the sample and results in the improvement of thermal shock resistance [15, 29]. Tamura et al. reported [29] that 1.5% nanocarbon batch showed equivalent thermal spalling properties to 18% graphite containing composition in MgO-C system. The batch
compositions with more than 2% nanocarbon showed less residual strength might be due to greater extent of oxidation of the agglomerated nanocarbon particles.
Conclusion: The batch compositions prepared with nanocarbon showed higher bulk density and cold crushing strength compared to the reference 25% graphite containing composition. The graphite content of the reference batch is replaced by fine alumina and nanocarbon in nanocarbon containing batches. Replacing lower density material (graphite) with higher density material (alumina) enhances density values. Because of its smaller size, nanocarbon can easily fill the gaps among the different sized aggregate fractions of the refractory thereby improves density. The strength values of the nanocarbon batches increases up to a coking temperature of 1200oC due to the formation of aluminum carbide phase but decreases at higher coking temperatures due to its dissociation. Nanocarbon is active due to its high surface area can initiate the reaction at much lower temperatures than the graphite flake. Oxidation resistance is measured in terms of weight loss, and the corresponding carbon loss is measured. The 25% graphite batch after decarburization due to oxidation does not retain any strength. Graphite flake has larger size compared to nanocarbon, which on oxidation leaves a bigger pore, does not allow the remaining alumina to sinter and resulting in weak structure. However, the nanocarbon batches show higher density in decarburized layer after oxidation due to the sintering of alumina particles in the composition. Dense decarburized layer will restrict further oxidation. The residual strength values are observed to decrease in all the compositions with the increase in number of thermal spalling cycles. The 1% nanocarbon-containing batch showed better retention of strength compared to all the other compositions. The increase in amount of nanocarbon content above 2% does not promote the properties any further. Acknowledgement: The authors thankfully acknowledge the financial support of CSIR, GOI for sponsored project, Grant-in-Aid, Ref No. 22(0767)/18/EMR-II. The authors also acknowledge the support of the staff of Department of Ceramic Engg, NIT, Rourkela for their help in experimentation; Mr. S. K. Shrivastava and Mr. B.C. Nanda of IFGL Refractories, India, for their support in sample preparation. Also special thanks to Mr. S. Chatterjee, Almatis, India for support in HMOR testing.
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Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: