- Email: [email protected]
Microstructures and mechanical properties of Al2O3–C refractories using nickel-loaded ultrafine microcrystalline graphite and silicon additives
Available online at www.sciencedirect.com
CERAMICS INTERNATIONAL
Ceramics International 40 (2014) 15783–15793 www.elsevier.com/locate/ceramint
Microstructures and mechanical properties of Al2O3–C refractories using nickel-loaded ultrafine microcrystalline graphite and silicon additives Heng Wanga, Yawei Lia,n, Shaobai Sanga, Shengli Jinb, Yibiao Xua, Kaibao Yangc, Shuzhong Yuc a
The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, PR China b Chair of Ceramics, Montanuniversität leoben, Leoben A-8700, Austria c Magang Refractory Co., Ltd., Maanshan Iron & Steel Co., Ltd., Maanshan 243000, PR China Received 11 July 2014; received in revised form 18 July 2014; accepted 20 July 2014 Available online 26 July 2014
Abstract Ultrafine microcrystalline graphite (UMCG) and nickel-loaded ultrafine microcrystalline graphite (NMCG) were introduced into Al2O3–C refractories to totally replace graphite flake (GF), respectively. Their interactions with silicon additive were exploited through the observation of microstructures, mechanical and thermo-mechanical properties of Al2O3–C refractories after firing at 800–1400 1C in a coke powder bed. It revealed that with NMCG, carbon nanotubes and SiC whiskers are in-situ formed out of nickel-catalytic pyrolysis of phenol resin binder at the temperature low to 1000 1C. Moreover, the amount of SiC whiskers increases with increasing fire temperature. The synergistic effect of the ultrafine microcrystalline graphite, carbon nanotubes and SiC whiskers contributes to the substantial improvement of mechanical properties and thermal shock resistance of Al2O3–C refractories using nickel-loaded ultrafine microcrystalline graphite as carbon source. & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved. Keywords: C. Thermal shock resistance; Al2O3–C refractories; Nickel-loaded ultrafine microcrystalline graphite; Carbon nanotubes; SiC whiskers
1. Introduction Carbon containing refractories have been widely used in the metallurgical industry due to their excellent mechanical properties, thermal shock resistance and corrosion resistance [1–3]. As a main carbon source, graphite flake (GF) is encountering the depletion of resources and consequently price increase. However, microcrystalline graphite (MCG) is abundant in nature and cheap which has attracted a lot of attention of researchers and producers of refractories recently [4]. Nowadays, the submicron or nano-carbon sources with large surface area, high activity, and high thermal conductivity were achieved by high-energy ball milling technique in carbon containing refractories [5]. In our previous work, ultrafine microcrystalline graphite (UMCG) prepared by high-energy ball milling technique was used to partly replace graphite flake n
Corresponding author. Tel./Fax: þ86 27 68862188. E-mail address: [email protected] (Y. Li).
http://dx.doi.org/10.1016/j.ceramint.2014.07.104 0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
in Al2O3–C refractories comprising of Al and Si additives, and brought about obviously positive strengthening effects [6]. Recently the microstructure design at nano-sized level for preparing high performance low carbon refractories with nanocarbon particles, nanooxide particles and organic binders has been put forward to modify traditional carbon containing refractories [7–18]. The applications of nanocarbon sources in the form of nano-carbon black (CB), carbon nanofibers (CNFs), carbon nanotubes (CNTs) and graphene or graphene oxide nanosheets (GONs) show their great potential in developing high-performance low carbon containing refractories. In the case of CNTs, Luo et al. [19] and Aneziris et al. [20] found that Al2O3–C refractories with the addition of CNTs showed better mechanical properties than those without CNTs. The achievement of superior properties originated from a combination of the nature of CNTs and in-situ formed ceramic whiskers in the matrix. Also, Roungos et al. [21]; reported that the addition of CNTs into Al2O3–C refractories gave rise to excellent thermal shock performance.
15784
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
The motivation of the present work is to form the carbon network with in-situ formed carbon nanotubes and SiC whiskers in the presence of silicon in Al2O3–C refractories. The following facts contribute to the above idea. The residual carbon of phenolic resin can be transformed into carbon nanotubes with nickel nitrate catalyst [22,23]. The morphology of in-situ ceramic phases in carbon containing refractories is in dependence of carbon sources. As shown in Ref. [24], the presence of graphite flake promotes the formation of SiC whiskers in Al2O3–C refractories. Similar to graphite flake, ultrafine microcrystalline graphite possesses the hexagonal structure besides its undeveloped micro-crystals structures. It is believed that the addition of UMCG also will benefit the formation of SiC whiskers. Therefore, nickel nitrate catalyst was doped into UMCGs by high-energy ball milling process and further introduced into Al2O3–C refractories combined with silicon additive.
2.2. Preparation of Al2O3–C refractories
2. Experimental
The raw materials used for preparing Al2O3–C specimens were tabular alumina (3–1 mm, 1–0.5 mm, 0.5–0 mm, r 0.045 mm and 20 μm, 98 wt% Al2O3, Zhejiang Zili Co., Ltd., China), silicon powder ( 45 μm, 98 wt% Si, China), graphite flake (r 0.074 mm, 97.58 wt% fixed carbon, Qingdao, China), as-prepared UMCG/α-Al2O3 and NMCG/α-Al2O3 composite powders. In addition, thermosetting phenolic resin (liquid, 440% fixed carbon, Wuhan Lifa Chemistry Industry Co., Ltd., China) was used as binder. The investigated Al2O3– C material ingredients with 2 wt% graphite content are presented in Table 1. All the raw materials were mixed for 30 min in a mixer with the rotating speed of 80–100 rpm. After kneading, the mixtures were pressed at 150 MPa into prismatic specimens (25 25 140 mm3) and then cured at 150 1C for 24 h. The cured specimens were put in a sagger filled with coke powder and fired at 800 1C, 1000 1C, 1200 1C and 1400 1C with a heating rate of 5 1C/min and holding time for 3 h.
2.1. Preparation of NMCG/α-Al2O3 and UMCG/α-Al2O3 composite powders
2.3. Testing and characterization methods
The nickel-loaded ultrafine microcrystalline graphite (NMCG) powders were prepared by directly mixing and dissolving Ni(NO3)2 6H2O (Analytical Reagent, Sinopharm, China) and absolute ethyl alcohol firstly; and then natural microcrystalline graphite (MCG, r 0.074 mm, 84.7 wt% fixed carbon, Chenzhou, China) was added into Ni(NO3)2 6H2O solution together with α-Al2O3 powders ( 2 μm, Al2O3 Z 99 wt%, Kaifeng Special Refractories Co., Ltd., China). The mixing mass ratio of MCG:Ni(NO3)2 6H2O :α-Al2O3 was 1:1:4. Finally the mixtures were wet-milled at a rotating rate of 400 rpm in a planetary balling for 7 h (corundum balls as the abrasive media) using absolute ethyl alcohol as the disperse media. The mass ratio of corundum balls to the powder mixtures was 2:1. After ball milling, the composite powders were dried at 70 1C for 24 h and then ground into fine powders. The ultrafine microcrystalline graphite (UMCG) powders were prepared by high-energy ball milling natural microcrystalline graphite and α-Al2O3 powders with the same ball milling process. The mixing mass ratio of MCG to α-Al2O3 was 1:5.
The apparent porosity (AP) and bulk density (BD) of all the fired specimens were determined according to Archimedes' Principle with kerosene as medium. Cold modulus of rupture (CMOR) and flexural modulus (FM) were measured by threepoint bending test at ambient temperature with a span of 80 mm and a loading rate of 0.5 mm/min by means of electronic digital control system (EDC120, DOLI Company, Germany). The force–displacement curve of each specimen was recorded simultaneously during the test. Three specimens of each material were used for the respective measurements. The phase compositions of the fired specimens were analyzed by X-ray diffraction (XRD, X’Pert Pro, Philips, Netherlands). The microstructures of all the fired Al2O3–C refractories were observed by a field emission scanning electron microscope (FESEM, Quanta 400, FEI Company, USA) equipped with energy dispersive X-ray spectroscope (EDS, phoenix, Philips, Eindhoven, The Netherlands) and high resolution transmission electron microscope (HRTEM, JEM-2100, JEOL, Japan). The thermal shock resistance of the specimens fired at 1400 1C was tested by firstly heating the specimens in a coke
Table 1 Ingredients of Al2O3–C refractories. Raw materials
S-GF
S-UMCG
S-NMCG
Tabular alumina α-Al2O3 Silicon Graphite flake UMCG/α-Al2O3 composite powders NMCG/α-Al2O3 composite powders Liquid phenolic resin
84 10 4 2 – – þ 4.5
84 – 4 – 12 – þ 4.5
84 – 4 – – 12 þ4.5
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
bed to 1100 1C with a heating rate of 5 1C/min and holding for 30 min, following it quick quenching of specimens into a water bath with room temperature was carried out. After one thermal shock cycle, the residual CMOR of the specimens was assessed by the three-point bending test. The residual strength ratio of CMOR was calculated by the change in CMOR before and after the thermal shocks, i.e. the residual strength ratio of CMOR ¼ 100 CMORTS/CMOR, where CMOR and CMORTS were the CMOR before and after one thermal shock cycle, respectively.
15785
microcrystalline graphite (UMCG) and nickel-loaded ultrafine microcrystalline graphite (NMCG) composite powders, respectively. Obviously, graphite flake and natural microcrystalline graphite were approximately 50 μm in length and 10 μm in thickness (Fig. 1a and b), while the ultrafine microcrystalline graphite and nickel-loaded ultrafine microcrystalline graphite with less than 5 μm in size (indicated by an arrow above) were distributed homogeneously in α-Al2O3 powders after the highenergy ball milling process, which were in agreement with that reported by the previous work [6], as shown in Fig. 1c and d, respectively.
3. Results and discussion 3.2. Phase composition of Al2O3–C refractories 3.1. Microstructure of different types of graphite Fig. 1a, b, c and d shows the SEM micrographs of graphite flake (GF), natural microcrystalline graphite (MCG), ultrafine
The phase composition of Al2O3–C refractories consisting of various carbon sources were investigated by means of XRD and shown in Fig. 2. In the case of conventional Al2O3–C refractories
Fig. 1. Microstructure of various carbon resources: (a) GF, (b) MCG, (c) UMCG and (d) NMCG.
15786
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
Fig. 2. XRD patterns of specimens S-GF (a), S-UMCG (b) and S-NMCG (c) fired at different temperatures.
containing graphite flake (Fig. 2a), no new phases were detected in the specimen after coked at 800–1200 1C, besides corundum, graphite and Si phases. After firing at 1400 1C, the diffraction peak of SiC appeared and the Si phase disappeared completely. With respect to the specimen containing ultrafine microcrystalline graphite, the phase transformation at 800 1C and 1000 1C was similar to those of the specimens S-GF. However, the SiC phase appeared in the specimen S-UMCG coked at 1200 1C and increased slightly at 1400 1C (Fig. 2b). When nickel-loaded ultrafine microcrystalline graphite was used, SiC phase was formed at 1000 1C and increased with increasing the temperatures from 1200 1C to 1400 1C (Fig. 2c). For specimens S-UMCG and S-NMCG, the peak intensity of graphite phase was lower than that of S-GF because the microcrystalline graphite had microcrystals structures and lower degree of graphitization. 3.3. Microstructures of Al2O3–C refractories The microstructure evolution of the Al2O3–C specimens containing various carbon resources fired at different
temperatures are presented in Figs. 3–5, respectively. It can be seen that graphite flake kept the initial microstructure and no new ceramic phases formed in the matrix of the specimen S-GF after firing at 800–1000 1C (Fig. 3a and b). However, after firing at 1200 1C, small amount of whiskers with nanosizes in diameter and micrometer size in length on the edge of the graphite could be observed in the specimen (Fig. 3c). Combing the EDS analysis and our previous researches, the whiskers are SiC phase [24]. It is worth noting that the diffraction peak of SiC was not detected at this temperature but the height of the peak of Si decreased, which might be attributed to the low amounts of SiC [25,26]. More SiC whiskers with larger length appeared in the specimen as the temperature increased to 1400 1C (Fig. 3d). Similarly, no new ceramic phases were detected in the specimen S-UMCG after firing at 800–1000 1C. However, the specimen obtained a relative dense structure owing to the ultrafine microcrystalline graphite dispersing homogeneously in the matrix (Fig. 4a and b). After firing at 1200 1C, a plenty of whiskers were found in the specimen S-UMCG, which were
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
15787
Fig. 3. SEM micrographs of fracture surfaces of Al2O3–C specimen S-GF fired at 800 1C (a), 1000 1C (b), 1200 1C (c), and 1400 1C (d).
confirmed to be SiC whiskers according to the XRD and EDS analysis (Fig. 4c). With the increase of firing temperature to 1400 1C, the SiC whiskers interlocked with each other to form intertextures and well distributed in the matrix (Fig. 4d). For the specimen S-NMCG fried at 800 1C, ultrafine microcrystalline graphite homogeneously dispersed in the matrix (Fig. 5a). After firing at 1000 1C, a large number of SiC whiskers and nanotubes with length less than 2 μm were well distributed in the matrix of the specimen (Fig. 5b). With increasing the firing temperature to 1200 1C, the quantity and size of the SiC whiskers and nanotubes increased and they also interlocked with each other (Fig. 5c). Compared with specimen S-UMCG, more SiC whiskers with larger length and diameter formed in the matrix of the specimen S-NMCG after firing at 1400 1C (Fig. 5d). In order to further investigate the microstructure of the SiC whiskers and nanotubes in the specimen S-NMCG fired at 1200 1C, the specimen was ground into powder and observed by HRTEM, as shown in Fig. 6. According to the HRTEM
analysis, the nanotubes are typical of carbon nanotubes (CNTs), which have a hollow structure of 26.57 nm in width (Fig. 6a and b). The high magnification (Fig. 6b) shows that the CNTs have about 50 graphite layers and are about 69.03 nm in diameter. Fig. 6d shows higher magnification image of the SiC nanowires (Fig. 6c), in which well-defined fringe separation of 0.26 nm is consistent with the d-spacing of (111) plane for β-SiC crystallographic planes [27], suggesting that the growth direction of the SiC nanowires was [111]. The microstructural evolutions of Al2O3–C refractories were in good agreement with the phase compositions. When the Al2O3–C refractories fired at high temperatures, the SiC whiskers could form by the main reactions (1)–(4) [19]. Compared with natural graphite flake, the planetary mill favored the cleavage of the particles and produced a local pressure about 2–6 GPa by high-energy ball milling process [28,29] to exfoliate natural microcrystalline graphite into ultrafine microcrystalline graphite with more defects, active
15788
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
Fig. 4. SEM micrographs of fracture surfaces of Al2O3–C specimen S-UMCG fired at 800 1C (a), 1000 1C (b), 1200 1C (c), and 1400 1C (d).
sites and higher activity to enhance the formation of the SiC whiskers [6,30] in the Al2O3–C refractories. Therefore, compared with the specimen S-GF, the diffraction peaks of SiC were detected at lower firing temperature (1200 1C) and more SiC whiskers formed in the specimen S-UMCG at the same heating temperature. Si (s, g)þ C (s)=SiC (s)
(1)
2Si (s, g)þ O2 (g)=2SiO (g)
(2)
SiO (g) þ 2C (s)=SiC (s) þ CO (g)
(3)
SiO (g) þ 3CO (g)=SiC (s) þ 2CO2 (g)
(4)
With regard to the specimen S-NMCG containing nickelloaded ultrafine microcrystalline graphite fired at high temperatures, firstly, the organic groups of phenolic resin would decompose into gaseous phases such as CH4, CO, CO2 and C2H6 during the heating process [31–33]. Then, the carbon
atoms from the gaseous phases diffusing towards and contacting with the catalyst nanoparticles form liquidized metal carbide. The continuous supply of the carbon source led to the over-saturation of carbon atoms in the nanoparticles, which would diffuse through metal particles to the other side of the crystal face to deposit and form CNTs [34,35] in the matrix. The growing mechanism of CNTs was generally regarded as “adsorption–diffusion–deposition” [36,37]. Moreover, in the specimen S-NMCG, the reaction between nickel and additive Si could promote the occurrence of gas–solid and gas–liquid– solid reaction mechanisms [38–40], which resulted in significant decrease in the formation temperature of SiC whiskers. Additionally, it has to be mentioned that the in-situ formed CNTs could also act as a kind of carbon source, which could accelerate the formation of SiC whiskers. Therefore, the SiC whiskers appeared at a much lower temperature and their quantity is highest in the specimen S-NMCG than other specimens at the same fired temperature.
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
15789
Fig. 5. SEM micrographs of fracture surfaces of Al2O3–C specimen S-NMCG fired at 800 1C (a), 1000 1C (b), 1200 1C (c), and 1400 1C (d).
3.4. Mechanical properties of Al2O3–C refractories Mechanical properties including CMOR and FM of Al2O3– C refractories specimens were tested by three-point bending method at room temperature. The results show that the firing temperature and types of carbon sources had an obvious influence on the CMOR and FM (Table 2). It can be found that the CMOR and FM of all the specimens increased with the increase of firing temperature and the specimens S-NMCG displayed highest value of CMOR and FM. The force–displacement curves of Al2O3–C refractories specimens fired at 800 1C, 1000 1C, 1200 1C and 1400 1C are depicted in Fig. 7. Obviously, the forces and displacements values of the specimens fired at different temperatures had the same variation tendency as the values of CMOR and FM. Namely, the forces and displacements increased continually with the increase of firing temperature from 800 1C to
1400 1C, indicating that the strength and toughness of the specimens increased with the elevated temperatures. Nevertheless, the specimen S-NMCG showed the largest displacement and biggest force at various firing temperatures. It is confirmed that Al2O3–C refractories specimens containing nickel-loaded ultrafine microcrystalline graphite possessed the best mechanical strength and toughness. The mechanical properties were closely associated with the physical properties such as apparent porosity and bulk density. Fig. 8 illustrates the evolutions of apparent porosity and bulk density of all specimens fired at various temperatures. The results show that the apparent porosity decreased with the temperature from 800 1C to 1200 1C, which were considered as a direct consequence of formation of new ceramic phases in the materials [41,42]. However, when the firing temperature reached to 1400 1C, considerable expansion occurred in the Al2O3–C specimens, resulting in increased of
15790
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
Fig. 6. TEM micrographs of specimen S-NMCG fired at 1200 1C. (a, b) CNTs and (c, d) SiC whiskers.
Table 2 CMOR and FM of all specimens fired at different temperatures. Temperature (1C)
Index
S-GF
S-UMCG
S-NMCG
800
CMOR (MPa) FM (GPa) CMOR (MPa) FM (GPa) CMOR (MPa) FM (GPa) CMOR (MPa) FM (GPa)
4.50 1.51 4.66 1.69 18.52 3.49 29.22 4.61
4.80 1.83 5.91 1.78 21.17 3.61 31.08 4.65
5.48 2.07 13.29 2.95 25.78 4.29 34.06 4.87
1000 1200 1400
the apparent porosity [19]. Correspondingly, the bulk density as a function of firing temperature was opposite to that of apparent porosity. In addition, it can be clearly seen that the specimens with ultrafine microcrystalline graphite or nickelloaded ultrafine microcrystalline graphite have lower apparent porosity and higher bulk density due to the fact that ultrafine
microcrystalline graphite could enhance the densification by filling the pore in the materials. The difference in microstructures and phase compositions unquestionably resulted in different mechanical properties of Al2O3–C refractories. The addition of nickel-loaded ultrafine microcrystalline graphite into Al2O3–C refractories presented
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
15791
Fig. 7. Force–displacement curves of all specimens fired at 800 1C (a), 1000 1C (b), 1200 1C (c), and 1400 1C (d).
Fig. 8. Apparent porosity (a) and bulk density (b) of all specimens fired at different temperatures.
the excellent mechanical properties among all the specimens. On the one hand, the lamellar nickel-loaded ultrafine microcrystalline graphite possessed a small particle size and high
specific surface area that could effectively fill the pore of the materials; also, similar to graphene oxide nanosheets, the ultrafine microcrystalline graphite powders had strengthening
15792
H. Wang et al. / Ceramics International 40 (2014) 15783–15793 Table 3 The CMOR change of all specimens fired at 1400 1C after one thermal shock cycle. Index
S-GF
S-UMCG
S-NMCG
CMOR (MPa) CMORTS (MPa) The residual strength ratio of CMOR (%)
29.22 7.07 24.09
31.08 8.74 28.12
34.06 12.41 36.44
and toughening mechanisms [6,43,44], including crack deflection, bridging, fracture and pull-out, strengthened the specimens. On the other hand, the in-situ formation of carbon nanotubes with excellent mechanical properties at 1000 1C, and presence of the nickel promoting formation of interlocking morphologies of SiC whiskers in the material at 1200– 1400 1C, lead to the superior mechanical properties of specimen S-NMCG. 3.5. Thermal shock resistance of Al2O3–C refractories The residual CMOR and residua strength ratio of CMOR of all specimens fired at 1400 1C after one thermal shock cycle are displayed in Table 3. As shown in the table, the specimen S-NMCG with nickel-loaded ultrafine microcrystalline graphite has the highest residual strength and residual strength ratio of CMOR, and the specimen S-UMCG is at the second rank. The residual strength ratio of specimen S-NMCG is 36.44% and decreases to 28.12% and 24.09% for the specimens S-UMCG and S-GF, respectively. The results indicate that the addition of nickel-loaded ultrafine microcrystalline graphite benefits the improvement of the thermal shock resistance of Al2O3–C refractories. It is clear that the homogeneously distributed micron and submicron ultrafine microcrystalline graphite powders and plenty of in-situ formed interlocking network ceramic whiskers together with the carbon nanotubes in the matrix contribute to the improvement of thermal shock resistance of Al2O3–C refractories [21,45]. 4. Conclusions The following conclusions can be made on the basis of phase compositions, microstructures mechanical properties and thermal shock resistance of Al2O3–C refractories containing nickel-loaded ultrafine microcrystalline graphite with silicon additive fired at 800–1400 1C in a coke powder bed. (1) The addition of nickel-loaded ultrafine microcrystalline graphite in Al2O3–C refractories can in-situ form carbon nanotubes out of catalytic pyrolysis of phenol resin binder as well as SiC whiskers at the temperature low to 1000 1C. Moreover, the amount of SiC whiskers increases with increasing fire temperature. (2) The mechanical properties and thermal shock resistance are greatly improved for the Al2O3–C refractories using nickel-loaded ultrafine microcrystalline graphite as carbon source in comparison with specimens containing ultrafine
microcrystalline graphite (UMCG) or graphite flake (GF), which can be explained by the synergistic effect of the ultrafine microcrystalline graphite, carbon nanotubes and SiC whiskers.
Acknowledgments The authors would like to thank the financial support from the National Natural Science Foundation of China (51372176), the Natural Science Foundation of Hubei Province (2013CFA105), Wuhan Science and Technology Bureau (2013010602010210 and 201271130449), and Yixing Science and Technology Bureau (20130419032713). References [1] E. Mohamed, M. Ewais, Carbon based refractories, J. Ceram. Soc. Jpn. 112 (2004) 517–532. [2] S. Zhang, W.E. Lee, Influence of additives on corrosion resistance and corroded microstructures of MgO–C refractories, J. Eur. Ceram. Soc. 21 (2001) 2393–2405. [3] C.G. Aneziris, U. Klippel, W. Schärfl, V. Stein, Y.W. Li, Functional refractory material design for advanced thermal shock performance due to titania additions, Int. J. Appl. Ceram. Technol. 4 (2007) 481–489. [4] X. Liang, Y.W. Li, Y. Zhang, S.B. Sang, Q.H. Wang, Effect of purified amorphous graphite on microstructure and mechanical properties of Al2O3–C refractories, J. Chin. Ceram. Soc. 42 (2014) 357–365. [5] X.X. Wu, H.X. Li, G.Q. Liu, C.C. Liu, G. Wang, J.L. Sun, Nanocarboncoated α-Al2O3 composite powders synthesized by high-energy ball milling, J. Inorg. Mater. 28 (2013) 261–266. [6] H. Wang, Y.W. Li, T.B. Zhu, S.B. Sang, Q.H. Wang, Microstructures and mechanical properties of Al2O3–C refractories with addition of microcrystalline graphite, Ceram. Int. 40 (2014) 11139–11148. [7] S. Tamura, T. Ochiai, S. Takanaga, T. Kanai, H. Nakamura, Nano-tech refractories-1: the development of nano structured matrix, in: Proceedings of the UNITECR’03 Congress, Osaka, Japan, 19–22 October, 2003. [8] S. Takanaga, T. Ochiai, S. Tamura, T. Kanai, H. Nakamura, Nano-tech. refractories-2: the application of the nano structural matrix to MgO–C bricks, in: Proceedings of the UNITECR’03 Congress, Osaka, Japan, 19–22 October, 2003. [9] S. Takanaga, Y. Fujiwara, M. Hatta, T. Ochiai, S..Tamura, Nano-tech. refractories-3: development of “MgO-rimmed MgO–C brick”, in: Proceedings of the UNITECR’05 Congress, Orlando, USA, 8–11 November, 2005. [10] Y. Shiratani, T. Yotabun, K. Chihara, T. Ochiai, S. Tamura, Nano-tech. refractories-4: the application of the nano structual matrix to SN plates, in: Proceedings of the UNITECR’05 Congress, Orlando, USA, 8–11 November, 2005. [11] M. Hatta, S. Takanaga, O. Matsuura, T. Ochiai, S. Tamura, Nano-tech. refractories-5: the application of B4C–C nano particles to MgO–C bricks, in: Proceedings of the UNITECR’07 Congress, Dresden, Germany, 18– 21 September, 2007.
H. Wang et al. / Ceramics International 40 (2014) 15783–15793 [12] S. Tamura, Y. Urushibara, O. Matsuura, T. Shin, Nano-tech. refractories6: observation of the texture after carbonization of nano-tech refractories, in: Proceedings of the UNITECR’07 Congress, Dresden, Germany, 18– 21 September, 2007. [13] H. Hattanda, T. Yotabun, T. Tsuda, T. Ochiai, S. Tamura, Nano-tech. refractories-7: application of nano structured matrix to SN plates, in: Proceedings of the UNITECR’07 Congress, Dresden, Germany, 18–21 September, 2007. [14] S. Tamura, T. Ochiai, S. Takanaga, T. Kanai, H. Nakamura, Nano-tech. refractories-8: technological philosophy and evolution of nano-tech. refractories, in: Proceedings of the UNITECR’11 Congress, Kyoto, Japan, October 30–November 2, 2011. [15] H. Yasumitsu, M. Hirashima, O. Matsuura, S. Takanaga, T. Ochiai, S. Tamura, Nano-tech. refractories-9: the basic study on the formation of the nano structured marix in MgO–C bricks, in: Proceedings of the UNITECR’11 Congress, Kyoto, Japan, October 30–November 2, 2011. [16] M. Tanaka, H. Kamio, J. Yoshitomi, T. Kayama, S. Hanagiri, K. Goto, Nano-tech. refractories-10: MgO–C bricks for converters to minimize the heat loss, in: Proceedings of the UNITECR’11 Congress, Kyoto, Japan, October 30–November 2, 2011. [17] H. Hattanda, T. Yotabun, T. Tsuda, H. Tanabe, Nano-tech. refractories11: the application of nano-technology to AG material for SN plates and carbon bricks for blast furnaces, in: Proceedings of the UNITECR’11 Congress, Kyoto, Japan, October 30–November 2, 2011. [18] S. Tamura, T. Ochiai, S. Takanaga, O. Matsuura, H. Yasumitsu, M. Hirashima, Nano-tech. refractories-12: development of MgO–C nano-tech. refractories of 0% graphite content, in: Proceedings of the UNITECR’13 Congress, Victoria, Canada, 10–13 September, 2013. [19] M. Luo, Y.W. Li, S.L. Jin, S.B. Sang, L. Zhao, Y.B. Li, Microstructures and mechanical properties of Al2O3–C refractories with addition of multiwalled carbon nanotubes, Mater. Sci. Eng. A 548 (2012) 134–141. [20] C.G. Aneziris, S.L. Jin, Y.W. Li, V. Roungos, Interactions of carbon nanotubes in Al2O3–C refractories for sliding gate applications, in: Proceedings of the UNTECR’09 Congress, Salvador, Brazil, 13–16 October, 2009. [21] V. Roungos, C.G. Aneziris, Improved thermal shock performance of Al2O3–C refractories due to nanoscaled additives, Ceram. Int. 38 (2012) 919–927. [22] M. Luo, Y.W. Li, S.B. Sang, L. Zhao, S.L. Jin, Y.B. Li, In situ formation of carbon nanotubes and ceramic whiskers in Al2O3–C refractories with addition of Ni-catalyzed phenolic resin, Mater. Sci. Eng. A 558 (2012) 533–542. [23] H. Wang, Y.W. Li, S.B. Sang, Q.H. Wang, Research on carbon nanotubes growth on the surface of expanded graphite by catalytic pyrolysis of the phenolic resin, Bull. Chin. Ceram. Soc. 33 (2014) 31–36. [24] H.B. Fan, Y.W. Li, S.B. Sang, Microstructures and mechanical properties of Al2O3–C refractories with silicon additive using different carbon sources, Mater. Sci. Eng. A 528 (2011) 3177–3185. [25] T.B. Zhu, Y.W. Li, M. Luo, S.B. Sang, Q.H. Wang, L. Zhao, Y.B. Li, S. J. Li, Microstructure and mechanical properties of MgO–C refractories containing graphite oxide nanosheets (GONs), Ceram. Int. 39 (2013) 3017–3025. [26] X.L. Chen, Y.W. Li, Y.B. Li, Effect of temperature on the properties and microstructures of carbon refractories for blast furnace (BF), Metall. Mater. Trans. A 40 (2009) 1675–1683. [27] G.Y. Li, X.D. Li, H. Wang, L. Liu, Ultra long SiC nanowires with fluctuating diameters synthesized in a polymer pyrolysis CVD route, Solid State Sci. 11 (2009) 2167–2172.
15793
[28] X.Q. Yue, H. Wang, S.Y. Wang, F.C. Zhang, R.J. Zhang, In-plane defects produced by ball-milling of expanded graphite, J. Alloys Compd. 505 (2010) 286–290. [29] X.Q. Yue, L. Li, R.J. Zhang, F.C. Zhang, Effect of expansion temperature of expandable graphite on microstructure evolution of expanded graphite during high-energy ball-milling, Mater. Charact. 60 (2009) 1541–1544. [30] Y.W. Li, Q.H. Wang, H.B. Fan, S.B. Sang, Y.B. Li, L. Zhao, Synthesis of silicon carbide whiskers using reactive graphite as template, Ceram. Int. 40 (2014) 1481–1488. [31] J.G. Wang, H.Y. Jiang, N. Jiang, Study on the pyrolysis of phenolformaldehyde (PF) resin and modified PF resin, Thermochim. Acta 496 (2009) 136–142. [32] J.G. Wang, N. Jiang, H.Y. Jiang, Micro-structural evolution of phenol– formal dehyde resin modified by boron carbide at elevated temperatures, Chem. Phys. 120 (2010) 187–192. [33] I. Stamatin, A. Morozan, A. Dumitru, V. Ciupina, G. Prodan, J. Niewolski, H. Figiel, The synthesis of multi-walled carbon nanotubes (MWNTs) by catalytic pyrolysis of the phenol–formaldehyde resins, Physica E 37 (2007) 44–48. [34] J.G. Zhao, Q.G. Guo, J.L. Shi, L. Liu, J.S. Jia, Y.C. Liu, H.Q. Wang, Carbon nanotube growth in the pores of expanded graphite by chemical vapor deposition, Carbon 47 (2009) 1747–1751. [35] J.G. Zhao, X.Y. Guo, Q.G. Guo, L. Gu, Y. Guo, F. Feng, Growth of carbon nanotubes on natural organic precursors by chemical vapor deposition, Carbon 49 (2011) 2155–2158. [36] J.G. Zhao, L. Liu, Q.G. Guo, J.L. Shi, G.T. Zhai, J.R. Song, Z.J. Liu, Growth of carbon nanotubes on the surface of carbon fibers, Carbon 46 (2008) 380–383. [37] J.R. Nielsen, D.L. Trimm, Mechanism of carbon formation on nickelcontaining catalysts, J. Catal. 48 (1977) 155–165. [38] R. Zhai, G.Y. Yang, R.B. Wu, J.J. Chen, J. Lin, L.L. Wu, Y. Pan, VLS growth of SiC whisker from FeSi flux, Acta Mater. Compos. Sin. 24 (2007) 97–102. [39] X.Y. Wu, G.Q. Jin, X.Y. Guo, Effect of the amounts of Fe catalyst on stacking faults and morphologies of β-SiC, New Carbon Mater. 20 (2005) 324–328. [40] F. Liang, N. Li, Effects of catalysts and calcinations temperatures on synthesis of SiC fibers, Refractories 43 (2009) 359–362. [41] X.X. Yi, Y.W. Li, S.B. Sang, B. Tong, H. Peng, J. Li, Effects of microsilica powder on microstructure and mechanical properties of fired alumina–zirconia–carbon sliding gate plates under different firing atmospheres, Refract. Worldforum 3 (2011) 107–112. [42] S. Zhang, N.J. Marriott, W.E. Lee, Thermo chemistry and microstructures of MgO–C refractories containing various antioxidants, J. Eur. Ceram. Soc. 21 (2001) 1037–1047. [43] S. Stankovich, D.A. Dikin, G.H.B. Dommett, K.M. Kohlhaas, Graphenebased composite materials, Nature 442 (2006) 282–286. [44] P. Kun, O. Tapaszto, F. Wéber, C. Balázsi, Determination of structural and mechanical properties of multilayer graphene added silicon nitridebased composites, Ceram. Int. 38 (2012) 211–216. [45] T.B. Zhu, Y.W. Li, S.L. Jin, S.B. Sang, Q.H. Wang, L. Zhao, Y.B. Li, S. J. Li, Microstructure and mechanical properties of MgO–C refractories containing expanded graphite, Ceram. Int. 39 (2013) 4529–4537.
CERAMICS INTERNATIONAL
Ceramics International 40 (2014) 15783–15793 www.elsevier.com/locate/ceramint
Microstructures and mechanical properties of Al2O3–C refractories using nickel-loaded ultrafine microcrystalline graphite and silicon additives Heng Wanga, Yawei Lia,n, Shaobai Sanga, Shengli Jinb, Yibiao Xua, Kaibao Yangc, Shuzhong Yuc a
The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, PR China b Chair of Ceramics, Montanuniversität leoben, Leoben A-8700, Austria c Magang Refractory Co., Ltd., Maanshan Iron & Steel Co., Ltd., Maanshan 243000, PR China Received 11 July 2014; received in revised form 18 July 2014; accepted 20 July 2014 Available online 26 July 2014
Abstract Ultrafine microcrystalline graphite (UMCG) and nickel-loaded ultrafine microcrystalline graphite (NMCG) were introduced into Al2O3–C refractories to totally replace graphite flake (GF), respectively. Their interactions with silicon additive were exploited through the observation of microstructures, mechanical and thermo-mechanical properties of Al2O3–C refractories after firing at 800–1400 1C in a coke powder bed. It revealed that with NMCG, carbon nanotubes and SiC whiskers are in-situ formed out of nickel-catalytic pyrolysis of phenol resin binder at the temperature low to 1000 1C. Moreover, the amount of SiC whiskers increases with increasing fire temperature. The synergistic effect of the ultrafine microcrystalline graphite, carbon nanotubes and SiC whiskers contributes to the substantial improvement of mechanical properties and thermal shock resistance of Al2O3–C refractories using nickel-loaded ultrafine microcrystalline graphite as carbon source. & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved. Keywords: C. Thermal shock resistance; Al2O3–C refractories; Nickel-loaded ultrafine microcrystalline graphite; Carbon nanotubes; SiC whiskers
1. Introduction Carbon containing refractories have been widely used in the metallurgical industry due to their excellent mechanical properties, thermal shock resistance and corrosion resistance [1–3]. As a main carbon source, graphite flake (GF) is encountering the depletion of resources and consequently price increase. However, microcrystalline graphite (MCG) is abundant in nature and cheap which has attracted a lot of attention of researchers and producers of refractories recently [4]. Nowadays, the submicron or nano-carbon sources with large surface area, high activity, and high thermal conductivity were achieved by high-energy ball milling technique in carbon containing refractories [5]. In our previous work, ultrafine microcrystalline graphite (UMCG) prepared by high-energy ball milling technique was used to partly replace graphite flake n
Corresponding author. Tel./Fax: þ86 27 68862188. E-mail address: [email protected] (Y. Li).
http://dx.doi.org/10.1016/j.ceramint.2014.07.104 0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
in Al2O3–C refractories comprising of Al and Si additives, and brought about obviously positive strengthening effects [6]. Recently the microstructure design at nano-sized level for preparing high performance low carbon refractories with nanocarbon particles, nanooxide particles and organic binders has been put forward to modify traditional carbon containing refractories [7–18]. The applications of nanocarbon sources in the form of nano-carbon black (CB), carbon nanofibers (CNFs), carbon nanotubes (CNTs) and graphene or graphene oxide nanosheets (GONs) show their great potential in developing high-performance low carbon containing refractories. In the case of CNTs, Luo et al. [19] and Aneziris et al. [20] found that Al2O3–C refractories with the addition of CNTs showed better mechanical properties than those without CNTs. The achievement of superior properties originated from a combination of the nature of CNTs and in-situ formed ceramic whiskers in the matrix. Also, Roungos et al. [21]; reported that the addition of CNTs into Al2O3–C refractories gave rise to excellent thermal shock performance.
15784
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
The motivation of the present work is to form the carbon network with in-situ formed carbon nanotubes and SiC whiskers in the presence of silicon in Al2O3–C refractories. The following facts contribute to the above idea. The residual carbon of phenolic resin can be transformed into carbon nanotubes with nickel nitrate catalyst [22,23]. The morphology of in-situ ceramic phases in carbon containing refractories is in dependence of carbon sources. As shown in Ref. [24], the presence of graphite flake promotes the formation of SiC whiskers in Al2O3–C refractories. Similar to graphite flake, ultrafine microcrystalline graphite possesses the hexagonal structure besides its undeveloped micro-crystals structures. It is believed that the addition of UMCG also will benefit the formation of SiC whiskers. Therefore, nickel nitrate catalyst was doped into UMCGs by high-energy ball milling process and further introduced into Al2O3–C refractories combined with silicon additive.
2.2. Preparation of Al2O3–C refractories
2. Experimental
The raw materials used for preparing Al2O3–C specimens were tabular alumina (3–1 mm, 1–0.5 mm, 0.5–0 mm, r 0.045 mm and 20 μm, 98 wt% Al2O3, Zhejiang Zili Co., Ltd., China), silicon powder ( 45 μm, 98 wt% Si, China), graphite flake (r 0.074 mm, 97.58 wt% fixed carbon, Qingdao, China), as-prepared UMCG/α-Al2O3 and NMCG/α-Al2O3 composite powders. In addition, thermosetting phenolic resin (liquid, 440% fixed carbon, Wuhan Lifa Chemistry Industry Co., Ltd., China) was used as binder. The investigated Al2O3– C material ingredients with 2 wt% graphite content are presented in Table 1. All the raw materials were mixed for 30 min in a mixer with the rotating speed of 80–100 rpm. After kneading, the mixtures were pressed at 150 MPa into prismatic specimens (25 25 140 mm3) and then cured at 150 1C for 24 h. The cured specimens were put in a sagger filled with coke powder and fired at 800 1C, 1000 1C, 1200 1C and 1400 1C with a heating rate of 5 1C/min and holding time for 3 h.
2.1. Preparation of NMCG/α-Al2O3 and UMCG/α-Al2O3 composite powders
2.3. Testing and characterization methods
The nickel-loaded ultrafine microcrystalline graphite (NMCG) powders were prepared by directly mixing and dissolving Ni(NO3)2 6H2O (Analytical Reagent, Sinopharm, China) and absolute ethyl alcohol firstly; and then natural microcrystalline graphite (MCG, r 0.074 mm, 84.7 wt% fixed carbon, Chenzhou, China) was added into Ni(NO3)2 6H2O solution together with α-Al2O3 powders ( 2 μm, Al2O3 Z 99 wt%, Kaifeng Special Refractories Co., Ltd., China). The mixing mass ratio of MCG:Ni(NO3)2 6H2O :α-Al2O3 was 1:1:4. Finally the mixtures were wet-milled at a rotating rate of 400 rpm in a planetary balling for 7 h (corundum balls as the abrasive media) using absolute ethyl alcohol as the disperse media. The mass ratio of corundum balls to the powder mixtures was 2:1. After ball milling, the composite powders were dried at 70 1C for 24 h and then ground into fine powders. The ultrafine microcrystalline graphite (UMCG) powders were prepared by high-energy ball milling natural microcrystalline graphite and α-Al2O3 powders with the same ball milling process. The mixing mass ratio of MCG to α-Al2O3 was 1:5.
The apparent porosity (AP) and bulk density (BD) of all the fired specimens were determined according to Archimedes' Principle with kerosene as medium. Cold modulus of rupture (CMOR) and flexural modulus (FM) were measured by threepoint bending test at ambient temperature with a span of 80 mm and a loading rate of 0.5 mm/min by means of electronic digital control system (EDC120, DOLI Company, Germany). The force–displacement curve of each specimen was recorded simultaneously during the test. Three specimens of each material were used for the respective measurements. The phase compositions of the fired specimens were analyzed by X-ray diffraction (XRD, X’Pert Pro, Philips, Netherlands). The microstructures of all the fired Al2O3–C refractories were observed by a field emission scanning electron microscope (FESEM, Quanta 400, FEI Company, USA) equipped with energy dispersive X-ray spectroscope (EDS, phoenix, Philips, Eindhoven, The Netherlands) and high resolution transmission electron microscope (HRTEM, JEM-2100, JEOL, Japan). The thermal shock resistance of the specimens fired at 1400 1C was tested by firstly heating the specimens in a coke
Table 1 Ingredients of Al2O3–C refractories. Raw materials
S-GF
S-UMCG
S-NMCG
Tabular alumina α-Al2O3 Silicon Graphite flake UMCG/α-Al2O3 composite powders NMCG/α-Al2O3 composite powders Liquid phenolic resin
84 10 4 2 – – þ 4.5
84 – 4 – 12 – þ 4.5
84 – 4 – – 12 þ4.5
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
bed to 1100 1C with a heating rate of 5 1C/min and holding for 30 min, following it quick quenching of specimens into a water bath with room temperature was carried out. After one thermal shock cycle, the residual CMOR of the specimens was assessed by the three-point bending test. The residual strength ratio of CMOR was calculated by the change in CMOR before and after the thermal shocks, i.e. the residual strength ratio of CMOR ¼ 100 CMORTS/CMOR, where CMOR and CMORTS were the CMOR before and after one thermal shock cycle, respectively.
15785
microcrystalline graphite (UMCG) and nickel-loaded ultrafine microcrystalline graphite (NMCG) composite powders, respectively. Obviously, graphite flake and natural microcrystalline graphite were approximately 50 μm in length and 10 μm in thickness (Fig. 1a and b), while the ultrafine microcrystalline graphite and nickel-loaded ultrafine microcrystalline graphite with less than 5 μm in size (indicated by an arrow above) were distributed homogeneously in α-Al2O3 powders after the highenergy ball milling process, which were in agreement with that reported by the previous work [6], as shown in Fig. 1c and d, respectively.
3. Results and discussion 3.2. Phase composition of Al2O3–C refractories 3.1. Microstructure of different types of graphite Fig. 1a, b, c and d shows the SEM micrographs of graphite flake (GF), natural microcrystalline graphite (MCG), ultrafine
The phase composition of Al2O3–C refractories consisting of various carbon sources were investigated by means of XRD and shown in Fig. 2. In the case of conventional Al2O3–C refractories
Fig. 1. Microstructure of various carbon resources: (a) GF, (b) MCG, (c) UMCG and (d) NMCG.
15786
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
Fig. 2. XRD patterns of specimens S-GF (a), S-UMCG (b) and S-NMCG (c) fired at different temperatures.
containing graphite flake (Fig. 2a), no new phases were detected in the specimen after coked at 800–1200 1C, besides corundum, graphite and Si phases. After firing at 1400 1C, the diffraction peak of SiC appeared and the Si phase disappeared completely. With respect to the specimen containing ultrafine microcrystalline graphite, the phase transformation at 800 1C and 1000 1C was similar to those of the specimens S-GF. However, the SiC phase appeared in the specimen S-UMCG coked at 1200 1C and increased slightly at 1400 1C (Fig. 2b). When nickel-loaded ultrafine microcrystalline graphite was used, SiC phase was formed at 1000 1C and increased with increasing the temperatures from 1200 1C to 1400 1C (Fig. 2c). For specimens S-UMCG and S-NMCG, the peak intensity of graphite phase was lower than that of S-GF because the microcrystalline graphite had microcrystals structures and lower degree of graphitization. 3.3. Microstructures of Al2O3–C refractories The microstructure evolution of the Al2O3–C specimens containing various carbon resources fired at different
temperatures are presented in Figs. 3–5, respectively. It can be seen that graphite flake kept the initial microstructure and no new ceramic phases formed in the matrix of the specimen S-GF after firing at 800–1000 1C (Fig. 3a and b). However, after firing at 1200 1C, small amount of whiskers with nanosizes in diameter and micrometer size in length on the edge of the graphite could be observed in the specimen (Fig. 3c). Combing the EDS analysis and our previous researches, the whiskers are SiC phase [24]. It is worth noting that the diffraction peak of SiC was not detected at this temperature but the height of the peak of Si decreased, which might be attributed to the low amounts of SiC [25,26]. More SiC whiskers with larger length appeared in the specimen as the temperature increased to 1400 1C (Fig. 3d). Similarly, no new ceramic phases were detected in the specimen S-UMCG after firing at 800–1000 1C. However, the specimen obtained a relative dense structure owing to the ultrafine microcrystalline graphite dispersing homogeneously in the matrix (Fig. 4a and b). After firing at 1200 1C, a plenty of whiskers were found in the specimen S-UMCG, which were
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
15787
Fig. 3. SEM micrographs of fracture surfaces of Al2O3–C specimen S-GF fired at 800 1C (a), 1000 1C (b), 1200 1C (c), and 1400 1C (d).
confirmed to be SiC whiskers according to the XRD and EDS analysis (Fig. 4c). With the increase of firing temperature to 1400 1C, the SiC whiskers interlocked with each other to form intertextures and well distributed in the matrix (Fig. 4d). For the specimen S-NMCG fried at 800 1C, ultrafine microcrystalline graphite homogeneously dispersed in the matrix (Fig. 5a). After firing at 1000 1C, a large number of SiC whiskers and nanotubes with length less than 2 μm were well distributed in the matrix of the specimen (Fig. 5b). With increasing the firing temperature to 1200 1C, the quantity and size of the SiC whiskers and nanotubes increased and they also interlocked with each other (Fig. 5c). Compared with specimen S-UMCG, more SiC whiskers with larger length and diameter formed in the matrix of the specimen S-NMCG after firing at 1400 1C (Fig. 5d). In order to further investigate the microstructure of the SiC whiskers and nanotubes in the specimen S-NMCG fired at 1200 1C, the specimen was ground into powder and observed by HRTEM, as shown in Fig. 6. According to the HRTEM
analysis, the nanotubes are typical of carbon nanotubes (CNTs), which have a hollow structure of 26.57 nm in width (Fig. 6a and b). The high magnification (Fig. 6b) shows that the CNTs have about 50 graphite layers and are about 69.03 nm in diameter. Fig. 6d shows higher magnification image of the SiC nanowires (Fig. 6c), in which well-defined fringe separation of 0.26 nm is consistent with the d-spacing of (111) plane for β-SiC crystallographic planes [27], suggesting that the growth direction of the SiC nanowires was [111]. The microstructural evolutions of Al2O3–C refractories were in good agreement with the phase compositions. When the Al2O3–C refractories fired at high temperatures, the SiC whiskers could form by the main reactions (1)–(4) [19]. Compared with natural graphite flake, the planetary mill favored the cleavage of the particles and produced a local pressure about 2–6 GPa by high-energy ball milling process [28,29] to exfoliate natural microcrystalline graphite into ultrafine microcrystalline graphite with more defects, active
15788
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
Fig. 4. SEM micrographs of fracture surfaces of Al2O3–C specimen S-UMCG fired at 800 1C (a), 1000 1C (b), 1200 1C (c), and 1400 1C (d).
sites and higher activity to enhance the formation of the SiC whiskers [6,30] in the Al2O3–C refractories. Therefore, compared with the specimen S-GF, the diffraction peaks of SiC were detected at lower firing temperature (1200 1C) and more SiC whiskers formed in the specimen S-UMCG at the same heating temperature. Si (s, g)þ C (s)=SiC (s)
(1)
2Si (s, g)þ O2 (g)=2SiO (g)
(2)
SiO (g) þ 2C (s)=SiC (s) þ CO (g)
(3)
SiO (g) þ 3CO (g)=SiC (s) þ 2CO2 (g)
(4)
With regard to the specimen S-NMCG containing nickelloaded ultrafine microcrystalline graphite fired at high temperatures, firstly, the organic groups of phenolic resin would decompose into gaseous phases such as CH4, CO, CO2 and C2H6 during the heating process [31–33]. Then, the carbon
atoms from the gaseous phases diffusing towards and contacting with the catalyst nanoparticles form liquidized metal carbide. The continuous supply of the carbon source led to the over-saturation of carbon atoms in the nanoparticles, which would diffuse through metal particles to the other side of the crystal face to deposit and form CNTs [34,35] in the matrix. The growing mechanism of CNTs was generally regarded as “adsorption–diffusion–deposition” [36,37]. Moreover, in the specimen S-NMCG, the reaction between nickel and additive Si could promote the occurrence of gas–solid and gas–liquid– solid reaction mechanisms [38–40], which resulted in significant decrease in the formation temperature of SiC whiskers. Additionally, it has to be mentioned that the in-situ formed CNTs could also act as a kind of carbon source, which could accelerate the formation of SiC whiskers. Therefore, the SiC whiskers appeared at a much lower temperature and their quantity is highest in the specimen S-NMCG than other specimens at the same fired temperature.
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
15789
Fig. 5. SEM micrographs of fracture surfaces of Al2O3–C specimen S-NMCG fired at 800 1C (a), 1000 1C (b), 1200 1C (c), and 1400 1C (d).
3.4. Mechanical properties of Al2O3–C refractories Mechanical properties including CMOR and FM of Al2O3– C refractories specimens were tested by three-point bending method at room temperature. The results show that the firing temperature and types of carbon sources had an obvious influence on the CMOR and FM (Table 2). It can be found that the CMOR and FM of all the specimens increased with the increase of firing temperature and the specimens S-NMCG displayed highest value of CMOR and FM. The force–displacement curves of Al2O3–C refractories specimens fired at 800 1C, 1000 1C, 1200 1C and 1400 1C are depicted in Fig. 7. Obviously, the forces and displacements values of the specimens fired at different temperatures had the same variation tendency as the values of CMOR and FM. Namely, the forces and displacements increased continually with the increase of firing temperature from 800 1C to
1400 1C, indicating that the strength and toughness of the specimens increased with the elevated temperatures. Nevertheless, the specimen S-NMCG showed the largest displacement and biggest force at various firing temperatures. It is confirmed that Al2O3–C refractories specimens containing nickel-loaded ultrafine microcrystalline graphite possessed the best mechanical strength and toughness. The mechanical properties were closely associated with the physical properties such as apparent porosity and bulk density. Fig. 8 illustrates the evolutions of apparent porosity and bulk density of all specimens fired at various temperatures. The results show that the apparent porosity decreased with the temperature from 800 1C to 1200 1C, which were considered as a direct consequence of formation of new ceramic phases in the materials [41,42]. However, when the firing temperature reached to 1400 1C, considerable expansion occurred in the Al2O3–C specimens, resulting in increased of
15790
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
Fig. 6. TEM micrographs of specimen S-NMCG fired at 1200 1C. (a, b) CNTs and (c, d) SiC whiskers.
Table 2 CMOR and FM of all specimens fired at different temperatures. Temperature (1C)
Index
S-GF
S-UMCG
S-NMCG
800
CMOR (MPa) FM (GPa) CMOR (MPa) FM (GPa) CMOR (MPa) FM (GPa) CMOR (MPa) FM (GPa)
4.50 1.51 4.66 1.69 18.52 3.49 29.22 4.61
4.80 1.83 5.91 1.78 21.17 3.61 31.08 4.65
5.48 2.07 13.29 2.95 25.78 4.29 34.06 4.87
1000 1200 1400
the apparent porosity [19]. Correspondingly, the bulk density as a function of firing temperature was opposite to that of apparent porosity. In addition, it can be clearly seen that the specimens with ultrafine microcrystalline graphite or nickelloaded ultrafine microcrystalline graphite have lower apparent porosity and higher bulk density due to the fact that ultrafine
microcrystalline graphite could enhance the densification by filling the pore in the materials. The difference in microstructures and phase compositions unquestionably resulted in different mechanical properties of Al2O3–C refractories. The addition of nickel-loaded ultrafine microcrystalline graphite into Al2O3–C refractories presented
H. Wang et al. / Ceramics International 40 (2014) 15783–15793
15791
Fig. 7. Force–displacement curves of all specimens fired at 800 1C (a), 1000 1C (b), 1200 1C (c), and 1400 1C (d).
Fig. 8. Apparent porosity (a) and bulk density (b) of all specimens fired at different temperatures.
the excellent mechanical properties among all the specimens. On the one hand, the lamellar nickel-loaded ultrafine microcrystalline graphite possessed a small particle size and high
specific surface area that could effectively fill the pore of the materials; also, similar to graphene oxide nanosheets, the ultrafine microcrystalline graphite powders had strengthening
15792
H. Wang et al. / Ceramics International 40 (2014) 15783–15793 Table 3 The CMOR change of all specimens fired at 1400 1C after one thermal shock cycle. Index
S-GF
S-UMCG
S-NMCG
CMOR (MPa) CMORTS (MPa) The residual strength ratio of CMOR (%)
29.22 7.07 24.09
31.08 8.74 28.12
34.06 12.41 36.44
and toughening mechanisms [6,43,44], including crack deflection, bridging, fracture and pull-out, strengthened the specimens. On the other hand, the in-situ formation of carbon nanotubes with excellent mechanical properties at 1000 1C, and presence of the nickel promoting formation of interlocking morphologies of SiC whiskers in the material at 1200– 1400 1C, lead to the superior mechanical properties of specimen S-NMCG. 3.5. Thermal shock resistance of Al2O3–C refractories The residual CMOR and residua strength ratio of CMOR of all specimens fired at 1400 1C after one thermal shock cycle are displayed in Table 3. As shown in the table, the specimen S-NMCG with nickel-loaded ultrafine microcrystalline graphite has the highest residual strength and residual strength ratio of CMOR, and the specimen S-UMCG is at the second rank. The residual strength ratio of specimen S-NMCG is 36.44% and decreases to 28.12% and 24.09% for the specimens S-UMCG and S-GF, respectively. The results indicate that the addition of nickel-loaded ultrafine microcrystalline graphite benefits the improvement of the thermal shock resistance of Al2O3–C refractories. It is clear that the homogeneously distributed micron and submicron ultrafine microcrystalline graphite powders and plenty of in-situ formed interlocking network ceramic whiskers together with the carbon nanotubes in the matrix contribute to the improvement of thermal shock resistance of Al2O3–C refractories [21,45]. 4. Conclusions The following conclusions can be made on the basis of phase compositions, microstructures mechanical properties and thermal shock resistance of Al2O3–C refractories containing nickel-loaded ultrafine microcrystalline graphite with silicon additive fired at 800–1400 1C in a coke powder bed. (1) The addition of nickel-loaded ultrafine microcrystalline graphite in Al2O3–C refractories can in-situ form carbon nanotubes out of catalytic pyrolysis of phenol resin binder as well as SiC whiskers at the temperature low to 1000 1C. Moreover, the amount of SiC whiskers increases with increasing fire temperature. (2) The mechanical properties and thermal shock resistance are greatly improved for the Al2O3–C refractories using nickel-loaded ultrafine microcrystalline graphite as carbon source in comparison with specimens containing ultrafine
microcrystalline graphite (UMCG) or graphite flake (GF), which can be explained by the synergistic effect of the ultrafine microcrystalline graphite, carbon nanotubes and SiC whiskers.
Acknowledgments The authors would like to thank the financial support from the National Natural Science Foundation of China (51372176), the Natural Science Foundation of Hubei Province (2013CFA105), Wuhan Science and Technology Bureau (2013010602010210 and 201271130449), and Yixing Science and Technology Bureau (20130419032713). References [1] E. Mohamed, M. Ewais, Carbon based refractories, J. Ceram. Soc. Jpn. 112 (2004) 517–532. [2] S. Zhang, W.E. Lee, Influence of additives on corrosion resistance and corroded microstructures of MgO–C refractories, J. Eur. Ceram. Soc. 21 (2001) 2393–2405. [3] C.G. Aneziris, U. Klippel, W. Schärfl, V. Stein, Y.W. Li, Functional refractory material design for advanced thermal shock performance due to titania additions, Int. J. Appl. Ceram. Technol. 4 (2007) 481–489. [4] X. Liang, Y.W. Li, Y. Zhang, S.B. Sang, Q.H. Wang, Effect of purified amorphous graphite on microstructure and mechanical properties of Al2O3–C refractories, J. Chin. Ceram. Soc. 42 (2014) 357–365. [5] X.X. Wu, H.X. Li, G.Q. Liu, C.C. Liu, G. Wang, J.L. Sun, Nanocarboncoated α-Al2O3 composite powders synthesized by high-energy ball milling, J. Inorg. Mater. 28 (2013) 261–266. [6] H. Wang, Y.W. Li, T.B. Zhu, S.B. Sang, Q.H. Wang, Microstructures and mechanical properties of Al2O3–C refractories with addition of microcrystalline graphite, Ceram. Int. 40 (2014) 11139–11148. [7] S. Tamura, T. Ochiai, S. Takanaga, T. Kanai, H. Nakamura, Nano-tech refractories-1: the development of nano structured matrix, in: Proceedings of the UNITECR’03 Congress, Osaka, Japan, 19–22 October, 2003. [8] S. Takanaga, T. Ochiai, S. Tamura, T. Kanai, H. Nakamura, Nano-tech. refractories-2: the application of the nano structural matrix to MgO–C bricks, in: Proceedings of the UNITECR’03 Congress, Osaka, Japan, 19–22 October, 2003. [9] S. Takanaga, Y. Fujiwara, M. Hatta, T. Ochiai, S..Tamura, Nano-tech. refractories-3: development of “MgO-rimmed MgO–C brick”, in: Proceedings of the UNITECR’05 Congress, Orlando, USA, 8–11 November, 2005. [10] Y. Shiratani, T. Yotabun, K. Chihara, T. Ochiai, S. Tamura, Nano-tech. refractories-4: the application of the nano structual matrix to SN plates, in: Proceedings of the UNITECR’05 Congress, Orlando, USA, 8–11 November, 2005. [11] M. Hatta, S. Takanaga, O. Matsuura, T. Ochiai, S. Tamura, Nano-tech. refractories-5: the application of B4C–C nano particles to MgO–C bricks, in: Proceedings of the UNITECR’07 Congress, Dresden, Germany, 18– 21 September, 2007.
H. Wang et al. / Ceramics International 40 (2014) 15783–15793 [12] S. Tamura, Y. Urushibara, O. Matsuura, T. Shin, Nano-tech. refractories6: observation of the texture after carbonization of nano-tech refractories, in: Proceedings of the UNITECR’07 Congress, Dresden, Germany, 18– 21 September, 2007. [13] H. Hattanda, T. Yotabun, T. Tsuda, T. Ochiai, S. Tamura, Nano-tech. refractories-7: application of nano structured matrix to SN plates, in: Proceedings of the UNITECR’07 Congress, Dresden, Germany, 18–21 September, 2007. [14] S. Tamura, T. Ochiai, S. Takanaga, T. Kanai, H. Nakamura, Nano-tech. refractories-8: technological philosophy and evolution of nano-tech. refractories, in: Proceedings of the UNITECR’11 Congress, Kyoto, Japan, October 30–November 2, 2011. [15] H. Yasumitsu, M. Hirashima, O. Matsuura, S. Takanaga, T. Ochiai, S. Tamura, Nano-tech. refractories-9: the basic study on the formation of the nano structured marix in MgO–C bricks, in: Proceedings of the UNITECR’11 Congress, Kyoto, Japan, October 30–November 2, 2011. [16] M. Tanaka, H. Kamio, J. Yoshitomi, T. Kayama, S. Hanagiri, K. Goto, Nano-tech. refractories-10: MgO–C bricks for converters to minimize the heat loss, in: Proceedings of the UNITECR’11 Congress, Kyoto, Japan, October 30–November 2, 2011. [17] H. Hattanda, T. Yotabun, T. Tsuda, H. Tanabe, Nano-tech. refractories11: the application of nano-technology to AG material for SN plates and carbon bricks for blast furnaces, in: Proceedings of the UNITECR’11 Congress, Kyoto, Japan, October 30–November 2, 2011. [18] S. Tamura, T. Ochiai, S. Takanaga, O. Matsuura, H. Yasumitsu, M. Hirashima, Nano-tech. refractories-12: development of MgO–C nano-tech. refractories of 0% graphite content, in: Proceedings of the UNITECR’13 Congress, Victoria, Canada, 10–13 September, 2013. [19] M. Luo, Y.W. Li, S.L. Jin, S.B. Sang, L. Zhao, Y.B. Li, Microstructures and mechanical properties of Al2O3–C refractories with addition of multiwalled carbon nanotubes, Mater. Sci. Eng. A 548 (2012) 134–141. [20] C.G. Aneziris, S.L. Jin, Y.W. Li, V. Roungos, Interactions of carbon nanotubes in Al2O3–C refractories for sliding gate applications, in: Proceedings of the UNTECR’09 Congress, Salvador, Brazil, 13–16 October, 2009. [21] V. Roungos, C.G. Aneziris, Improved thermal shock performance of Al2O3–C refractories due to nanoscaled additives, Ceram. Int. 38 (2012) 919–927. [22] M. Luo, Y.W. Li, S.B. Sang, L. Zhao, S.L. Jin, Y.B. Li, In situ formation of carbon nanotubes and ceramic whiskers in Al2O3–C refractories with addition of Ni-catalyzed phenolic resin, Mater. Sci. Eng. A 558 (2012) 533–542. [23] H. Wang, Y.W. Li, S.B. Sang, Q.H. Wang, Research on carbon nanotubes growth on the surface of expanded graphite by catalytic pyrolysis of the phenolic resin, Bull. Chin. Ceram. Soc. 33 (2014) 31–36. [24] H.B. Fan, Y.W. Li, S.B. Sang, Microstructures and mechanical properties of Al2O3–C refractories with silicon additive using different carbon sources, Mater. Sci. Eng. A 528 (2011) 3177–3185. [25] T.B. Zhu, Y.W. Li, M. Luo, S.B. Sang, Q.H. Wang, L. Zhao, Y.B. Li, S. J. Li, Microstructure and mechanical properties of MgO–C refractories containing graphite oxide nanosheets (GONs), Ceram. Int. 39 (2013) 3017–3025. [26] X.L. Chen, Y.W. Li, Y.B. Li, Effect of temperature on the properties and microstructures of carbon refractories for blast furnace (BF), Metall. Mater. Trans. A 40 (2009) 1675–1683. [27] G.Y. Li, X.D. Li, H. Wang, L. Liu, Ultra long SiC nanowires with fluctuating diameters synthesized in a polymer pyrolysis CVD route, Solid State Sci. 11 (2009) 2167–2172.
15793
[28] X.Q. Yue, H. Wang, S.Y. Wang, F.C. Zhang, R.J. Zhang, In-plane defects produced by ball-milling of expanded graphite, J. Alloys Compd. 505 (2010) 286–290. [29] X.Q. Yue, L. Li, R.J. Zhang, F.C. Zhang, Effect of expansion temperature of expandable graphite on microstructure evolution of expanded graphite during high-energy ball-milling, Mater. Charact. 60 (2009) 1541–1544. [30] Y.W. Li, Q.H. Wang, H.B. Fan, S.B. Sang, Y.B. Li, L. Zhao, Synthesis of silicon carbide whiskers using reactive graphite as template, Ceram. Int. 40 (2014) 1481–1488. [31] J.G. Wang, H.Y. Jiang, N. Jiang, Study on the pyrolysis of phenolformaldehyde (PF) resin and modified PF resin, Thermochim. Acta 496 (2009) 136–142. [32] J.G. Wang, N. Jiang, H.Y. Jiang, Micro-structural evolution of phenol– formal dehyde resin modified by boron carbide at elevated temperatures, Chem. Phys. 120 (2010) 187–192. [33] I. Stamatin, A. Morozan, A. Dumitru, V. Ciupina, G. Prodan, J. Niewolski, H. Figiel, The synthesis of multi-walled carbon nanotubes (MWNTs) by catalytic pyrolysis of the phenol–formaldehyde resins, Physica E 37 (2007) 44–48. [34] J.G. Zhao, Q.G. Guo, J.L. Shi, L. Liu, J.S. Jia, Y.C. Liu, H.Q. Wang, Carbon nanotube growth in the pores of expanded graphite by chemical vapor deposition, Carbon 47 (2009) 1747–1751. [35] J.G. Zhao, X.Y. Guo, Q.G. Guo, L. Gu, Y. Guo, F. Feng, Growth of carbon nanotubes on natural organic precursors by chemical vapor deposition, Carbon 49 (2011) 2155–2158. [36] J.G. Zhao, L. Liu, Q.G. Guo, J.L. Shi, G.T. Zhai, J.R. Song, Z.J. Liu, Growth of carbon nanotubes on the surface of carbon fibers, Carbon 46 (2008) 380–383. [37] J.R. Nielsen, D.L. Trimm, Mechanism of carbon formation on nickelcontaining catalysts, J. Catal. 48 (1977) 155–165. [38] R. Zhai, G.Y. Yang, R.B. Wu, J.J. Chen, J. Lin, L.L. Wu, Y. Pan, VLS growth of SiC whisker from FeSi flux, Acta Mater. Compos. Sin. 24 (2007) 97–102. [39] X.Y. Wu, G.Q. Jin, X.Y. Guo, Effect of the amounts of Fe catalyst on stacking faults and morphologies of β-SiC, New Carbon Mater. 20 (2005) 324–328. [40] F. Liang, N. Li, Effects of catalysts and calcinations temperatures on synthesis of SiC fibers, Refractories 43 (2009) 359–362. [41] X.X. Yi, Y.W. Li, S.B. Sang, B. Tong, H. Peng, J. Li, Effects of microsilica powder on microstructure and mechanical properties of fired alumina–zirconia–carbon sliding gate plates under different firing atmospheres, Refract. Worldforum 3 (2011) 107–112. [42] S. Zhang, N.J. Marriott, W.E. Lee, Thermo chemistry and microstructures of MgO–C refractories containing various antioxidants, J. Eur. Ceram. Soc. 21 (2001) 1037–1047. [43] S. Stankovich, D.A. Dikin, G.H.B. Dommett, K.M. Kohlhaas, Graphenebased composite materials, Nature 442 (2006) 282–286. [44] P. Kun, O. Tapaszto, F. Wéber, C. Balázsi, Determination of structural and mechanical properties of multilayer graphene added silicon nitridebased composites, Ceram. Int. 38 (2012) 211–216. [45] T.B. Zhu, Y.W. Li, S.L. Jin, S.B. Sang, Q.H. Wang, L. Zhao, Y.B. Li, S. J. Li, Microstructure and mechanical properties of MgO–C refractories containing expanded graphite, Ceram. Int. 39 (2013) 4529–4537.