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TiC composite fabricated by reactive melt infiltration
Accepted Manuscript Structural design of laminated B4C/TiC composite fabricated by reactive melt infiltration Mengyong Sun, Yuhang Bai, Mingxing Li, Shangwu Fan, Laifei Cheng PII:
S0925-8388(18)32392-2
DOI:
10.1016/j.jallcom.2018.06.271
Reference:
JALCOM 46607
To appear in:
Journal of Alloys and Compounds
Received Date: 1 February 2018 Revised Date:
13 June 2018
Accepted Date: 23 June 2018
Please cite this article as: M. Sun, Y. Bai, M. Li, S. Fan, L. Cheng, Structural design of laminated B4C/ TiC composite fabricated by reactive melt infiltration, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.06.271. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. 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.
ACCEPTED MANUSCRIPT Structural design of laminated B4C/TiC composite fabricated by reactive melt infiltration Mengyong Sun, Yuhang Bai, Mingxing Li, Shangwu Fan*, Laifei Cheng*
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Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University, Xi’an, Shaanxi 710072, PR China Abstract:
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Laminated boron carbide/titanium carbide (B4C/TiC) ceramic with multilayer structural unit was fabricated by reactive melt infiltration (RMI) process with silicon
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(Si) and aluminum silicon (Al40Si60) alloy. Laminated B4C/TiC ceramics infiltrated at a low temperature of 1300 °C with Al40Si60 alloy acquired structural unit including two layers of Al3B48C2–SiC, one layer of TiC–Ti3Si(Al)C2, and two layers of in-situ
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synthesized TiB2 at the interface between Al3B48C2 and TiC layers. The thickness of Al3B48C2–SiC layer, TiB2–SiC layer, and TiC–Ti3Si(Al)C2–SiC layer was found to be about 450, 30, and 90 µm, respectively. Reaction products of Al and B4C consisted of
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two types of Al–B–C ternary phases including Al3B48C2 and AlB24C4 attributed to the
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solution of Al in B4C and diffusion of Al into B4C. The diffusion of elements caused by concentration gradient and different diffusion rate of Si led to the formation of TiSi2 layer at 1550 °C and TiB2 layer at 1300 °C between B4C and TiC layers. Keywords: Structural design; Laminated ceramics; In-suit formation; Boron carbide– aluminum
1 *Corresponding author: Laifei Cheng, Ph.D., Professor; E-mail address: [email protected]. Shangwu Fan, Professor; E-mail address: [email protected].
ACCEPTED MANUSCRIPT 1. Introduction Boron carbide (B4C) is high-strength and low-weight ceramic with high hardness in combination with high wear resistance, good impact resistance, high melting point,
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resistance to chemical attack, and good electrical conductivity [1-6]. However, the dense B4C ceramics with low fracture toughness (about 3.7 MPa·m1/2) are usually fabricated by hot pressing (HP) and spark plasma sintering (SPS) at high temperature
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(above 2000 °C) and pressure (30 MPa), which limit their potential application as
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structural materials due to undesirable sample shape and expensive product [7-9]. Therefore, in order to overcome the above mentioned limitations, extensive research efforts have been devoted to the investigation of the pressureless sintering of B4C to fabricate ceramics with large size and complex shape. In pressureless sintering, B4C
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ceramics can also be densified using fine grain powders (<3 µm) and high sintering temperature (above 2150 °C) [10, 11]. In order to enhance the sinterability and mechanical properties of B4C ceramics by pressureless sintering, the addition of
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secondary phases (such as C, B, Si, Al, and Mg) as sintering aids has been carried out
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to decrease the sintering temperature of B4C ceramics [12-15]. Moreover, the reactive melt infiltration (RMI) is a rapid and economical pressureless sintering method to fabricate B4C ceramics by infiltrating molten silicon (silicon or aluminum silicon alloy) spontaneously into porous preforms under vacuum at a low temperature (below 1500 °C) [16-20]. Advantages of RMI include short fabrication time, cost effectiveness, and fabrication of near net shape ceramics. Therefore, understanding the phase equilibrium of the Si or alloy–B4C composites during the permeation
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molten Si into porous preforms at the temperature of 1450, 1650, and 1750 °C. They found that the B12(B, C, Si)3 phase precipitated at the defective concave surface of the large B4C grains. Frage et al. [23] reported that the Al–Si alloys with Si content higher
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than 40 wt.% could prevent the formation of brittle compound Al4C3. Viala et al. [24,
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25] systematically studied the chemical reactivity of B4C with Al during the RMI process. They found B4C to be unstable and it reacted with either solid or molten Al. When the temperature was above 868 °C, the reaction products were Al3BC and Al3B48C2 (β-AlB12) phases. Lee et al. [26] found numerous reactant phases including
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AlB2, Al3B48C2, AlB24C4 (AlB10), and Al3BC phases at the B4C–Al boundaries by the pressureless infiltration technique at 800 °C for 1 h. Arslan et al. [8] reported that the infiltration process should be performed at higher temperature of 1200 °C due to the
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wettability requirement between Al and B4C.
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Furthermore, layered structure design of ceramics is also an effective way to further improve the fracture toughness [27-31]. The laminated ceramic composites offer the best prospect for rational use of the unique physical and mechanical properties of monolithic ceramics, thus providing a way to improve the durability; toughness; and wear, corrosion, and thermal resistance. Recently, the MAX phase with nano-layered structure has been used to improve the fracture toughness [32-34]. Wang et al. investigated the Ti3Si(Al)C2-based ceramics synthesized by RMI of
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ACCEPTED MANUSCRIPT Al70Si30 alloy into the porous TiC preforms [35]. Therefore, the low temperature preparation by RMI process and multi-scale layer structure design can be combined for the fabrication of the laminated B4C ceramics with high toughness at low cost.
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In this study, the laminated B4C/titanium carbide (B4C/TiC) ceramics were fabricated by tape casting, lamination, and RMI process with silicon (Si) and aluminum silicon (Al40Si60) alloy. The structural design of laminated ceramics was
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investigated to study the interfacial reaction and phase transformation during the RMI
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process. Finally, the microstructures and mechanical properties of the laminated B4C/TiC composite were characterized. 2. Experiments
Tape casting is the most common method to fabricate the laminated ceramics.
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Fig. 1 shows the fabrication process of green body. For tape casting, slurries of compound powders (50 wt.%) were made by mixing powder blends with solvent (mixed alcohol and butanone, 40 wt.%) and dispersing agent (triethylphosphate, 2
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wt.%), and then milled for 24 h. Subsequently, binder (polyvinyl butyral (PVB), 4
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wt.%), plasticizer (dioctyl phthalate (DOP), 2 wt.%), and glycerol (2 wt.%) were added and the slurries were milled for an additional 24 h. Further, the slurries were degassed, and tape casting was carried out using an experimental tape casting equipment (Beijing Orient Sun-Tec Co., Ltd.) with a dual doctor blade casting head. After drying, the B4C–SiC green tape including B4C (75 wt.%, D50 = 20 µm), SiC (20 wt.%, D50 = 20 µm), and graphite powders (5 wt.%, D50 = 1 µm); and the TiC green tape including TiC (D50 = 1 µm) and graphite (10 wt.%) were fabricated. The
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layers of B4C–SiC green tapes and one layer of TiC green tape alternately.
Fig. 1. Schematic illustration of the fabrication process of green body.
Further, the green compact was laminated by low thermal compression at 150 °C
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under a pressure of 30 MPa. Then, the binder removal process was performed at 600 °C for 3 h with a heating rate of 1 °C min−1 in argon flow followed by sintering at
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1300 °C with a holding time of 1 h. Finally, the laminated B4C/TiC compacts with porosity of about 40% were infiltrated under vacuum (10 Pa) at 1550 °C for 1 h with
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Si designated as sample S1 and at 1300°C for 2 h with Al40Si60 alloy designated as sample S2, respectively. Open porosity was measured by the Archimedes’ method. Microstructures of the
samples were systematically characterized by scanning electron microscopy (SEM, S-4700, Japan) and optical microscopy (OM, OLYMPUS, Japan). The elemental composition and spatial distribution of ceramics were characterized by energy dispersive X-ray spectroscopy (EDS). The phase composition of materials was 5
ACCEPTED MANUSCRIPT characterized by the X-ray diffraction (XRD). The bending strength was measured via a three-point bending test using an electromechanical universal testing machine (CMT5504, MTS SYSTEMS, Co., Ltd., China). Fracture toughness (KIC) was
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evaluated by single-edge notched beam (SENB) test. The Vickers hardness was measured by Vickers’ indentation with a 9.8 N load applied for 15 s on polished
the FactSageTM software package (version 5.3) [36].
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3. Results and discussion
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sections (AMH43, LECO, USA). Thermodynamic analysis was carried out by using
The laminated B4C/TiC green bodies were fabricated by RMI process with Si at around 1550 °C and Al40Si60 alloy at 1300 °C. The siliconized sample S1 infiltrated at 1550 °C using Si exhibited large deformation, such as cracks and delamination,
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during the process of cooling, which resulted in fragmentation of the sample. The siliconized samples S2 infiltrated at 1300 °C using Al40Si60 alloy exhibited nearly dense composites with the porosity of 2 %. During the process of infiltration, the
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possible reactions are as follows [32, 37-41]: 3TiC(s) + 2Si(l) → Ti 3SiC2 (s) + SiC(s)
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(1)
TiC(s) + 3Si(l) → TiSi 2 (s) + SiC(s)
(2)
TiSi 2 (s) + Al-Si(l) → Ti-Si-Al(l)
(3)
2TiC + Ti-Si-Al(l) → Ti3Si(Al)C2 (s)
(4)
C(s) + Si(l) → SiC(s)
(5)
B4C(s) + 2 TiC(s) + 3Si(l) → 2TiB2 (s) + 3SiC(s)
(6)
Al(l) + B4 C(s) → Al3 B48C2 (s) / Al B24C4 (s)
(7)
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°C.
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Fig. 2. XRD patterns of laminated B4C/TiC ceramics fabricated by RMI process using Si at 1500
Fig. 3. XRD patterns of laminated B4C/TiC ceramics fabricated by RMI process using Al40Si60 alloy at 1300 °C.
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Figs. 2 and 3 exhibit the XRD patterns of samples S1 and S2, respectively. The
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XRD analysis results revealed that Ti3Si(Al)C2, Ti3SiC2, Al3B48C2, TiB2, and SiC were formed by RMI process. The residual Si, Al, and TiC were also detected. TiSi2 and Ti3SiC2 were formed by the reactions between TiC particles and Si (Eqs. 1 and 2). For the formation of Ti3Si(Al)C2, the ternary Ti–Si–Al melt was formed first, attributed to the dissolution of TiSi2 into Al–Si melt (Eq. 3). The TiC twins were formed due to decreasing boundary energy in the presence of Al, and then the Ti3SiC2 grains nucleated from TiC twins. Further, Al dissolved into Ti3SiC2 to form Ti3Si(Al)C2 (Eq.
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B4C (Eq. 6).
Fig. 4. Optical microscopy images (a) and (b) of sample S1 infiltrated at 1550 °C with Si and (c) and (d) of sample S2 infiltrated with Al40Si60 alloy at 1300 °C.
Fig. 4 shows the optical micrographs of the laminated B4C/TiC ceramics
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fabricated at different temperatures. The thickness ratio of B4C layer to TiC layer is
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about 3:1. Large size particles of B4C and SiC evenly disperse in the thick layer, and Ti3SiC2 (sample S1) or Ti3Si(Al)C2 (sample S2) locates in the thin layer. Figs. 4(b) and 4(d) show that TiSi2 and TiB2 layers were in-situ formed at the interface between B4C and TiC layers of samples S1 and S2, respectively. Fig. 4(a) exhibits large cracks in the sample S1 generated from the process of cooling, which resulted in breakdown of ceramic into small pieces. As a result, large size of laminated B4C/TiC ceramics
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no cracks are observed in the sample S2 filled with solid Si and Al.
Fig. 5. BSE images of sample S1 infiltrated with Si at 1550 °C
Fig. 5 shows the back-scattered electron (BSE) images of sample S1 and Table 1
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lists the EDS results of the points marked in Fig. 5, clearly indicating the distribution of TiSi2, Ti3SiC2, and SiC phases. In the B4C–SiC layer, the SiC and B4C particles are
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encapsulated by solid Si as shown in Fig. 5(b). Besides, a small amount of B12(C, Si, B)3 was also formed in the liquid Si environment based on EDS result of point 6 as
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shown in Fig. 5(b) [22]. Figs. 5(c) and 5(d) exhibit that the in-situ formed TiSi2 phase is located at the interface between B4C and TiC layers, and also inside the TiC layer. Further, Fig. 5(d) demonstrates almost complete conversion of all TiC phases into Ti3SiC2, TiSi2, and small-sized SiC phases. Formation of TiSi2 layer between B4C layer and TiC layer is attributed to the diffusion of elements caused by concentration gradient, and tensile stress due to the volume shrinkage at the interface [42]. The high temperature of 1550 °C (above the melt point of Si) led to faster diffusion of Si into 9
ACCEPTED MANUSCRIPT the interface than B, which resulted in the formation of TiSi2 (Eq. 2). Table 1 EDS results of the point marked in Fig. 5. 2
3
4
5
6
7
40.21 16.45 43.34
72.67 27.33
45.44 02.17 52.39
53.51 46.49 -
72.05 27.59 0.35 -
45.62 13.83 40.55 -
100.00 -
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B C Si Ti
1
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Point Element (at. %)
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Fig. 6 BSE images of sample S2 infiltrated at 1300 °C with Al40Si60 alloy.
Fig. 6 shows the BSE images of sample S2 infiltrated at 1300 °C with Al40Si60
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alloy. The laminated structure of B4C/TiC ceramic is composed of Al3B48C2–SiC–Si, TiB2–SiC, and TiC–Ti3Si(Al)C2–SiC phases. The thickness of the Al3B48C2–SiC–Si layer, TiB2–SiC layer, and TiC–Ti3Si(Al)C2–SiC layer is about 450, 30, and 90 µm, respectively. The in-situ formed TiB2 layer is located in the interface between B4C and TiC layers, which is attributed to the diffusion of Ti and B into the interface due to concentration gradient in Si–Al melt. Owing to the diffusion and reaction of B4C, the SiC was left on one side of the TiB2 layer. Apparent growth of grains could also be 10
ACCEPTED MANUSCRIPT seen, which is attributed to the enough holding time at 1300 °C for 2 h and the suitable liquid phase environment. Formation of TiB2 layer between TiC layer and B4C layer rather than TiSi2 could be attributed to low diffusion rate of Si caused by
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low temperature of 1300 °C and interfaces of Al blocking mass transfer. Consequently, the concentration gradient of Ti and B made the elements diffuse into the interface, which led to the formation of TiB2 (Eq. 6).
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Fig. 7 shows the EDS mapping results and BSE images of sample S2. Fig. 7(a)
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demonstrates that the residual Al and Si are independent rather than Al–Si alloy, and are mainly located between large size SiC particles. The luminance decreases with decreasing content of Si in Si and SiC phases. Fig. 7(b) shows that in the TiC layer, the phase consists mainly of Ti3Si(Al)C2, a small amount of TiSi2, residual TiC, and
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dendritic-like structure of SiC. In the B4C layer as shown in Fig. 7(c), the phase consists mainly of SiC phase and two types of Al–B–C phases resulting from the Eq. 7. According to the EDS results listed in Table 2, the gray area represents the Al3B48C2
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phase with the high Al content, and the dark gray area denotes the AlB24C4 phase with
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the low Al content. Owing to the heat treatment at 1300 °C and sufficient holding time, the reaction products of Al and B4C include mainly Al3B48C2 phase. Besides, the reaction could be regarded as the solution and diffusion of Al into B4C, thus the AlB24C4 phase was mainly found to be located in the interior of the Al–B–C phase.
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Fig. 7. (a) EDS mapping results of Al and Si and (b), (c) BSE images of sample S2 infiltrated
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at 1300 °C with Al40Si60 alloy. Table 2 EDS results of the point marked in Figs. 6 and 7. Point
1
Element (at.%)
13.18 86.82 -
3
4
48.70 5.99 40.11 5.20
39.66 9.95 10.42 3.73 36.23
47.54 3.72 48.74
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B C Al Si Ti
2
5
6
7
30.13 5.39 17.34 47.14
60.11 38.46 1.43 -
73.06 20.29 6.65 -
Fig. 8 shows crack propagation and fracture surface of sample S2. Fig. 8(a)
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demonstrates that the crack is obviously deflected at the interface between the B4C and the TiC layers, which was due to the residual stress caused by different thermal
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expansion coefficients. Figs. 8(b) and 8(c) show that the mode of strong binding occurs at the interface, and the fracture modes include transgranular fracture, and intergranular-fracture. Fig. 8(d) exhibits the distribution of Ti3Si(Al)C2 phase in the TiC layer.
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Fig. 8. (a) Micrographs of crack propagation and (b), (c), and (d) fracture surface of sample S2
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infiltrated at 1300 °C with Al40Si60 alloy.
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Fig. 9. Load-displacement curves of sample S2 infiltrated at 1300 °C with Al40Si60 alloy.
Fig. 9 shows the load-displacement curves with the step-like fracture mode for
sample S2, which could be attributed to the laminated structure with compressive layers. In the laminated ceramics with strong interface, the crack can be deflected by residual compressive stress at the interface, which could improve the toughness. The fracture toughness increases with the increase in the crack length, and this is called R-curve behavior [43]. The flexural strength and the fracture toughness of laminated
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ACCEPTED MANUSCRIPT Al3B48C2–SiC/TiC–Ti3Si(Al)C2 ceramics were 160 ± 2 MPa and 5 ± 0.3 MPa·m1/2, respectively. The Al3B48C2–SiC and TiC–Ti3Si(Al)C2 layers exhibited the Vickers’ hardness of 27.5 ± 0.5 and 15 ± 1 GPa, respectively.
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Fig. 10 exhibits the schematic depiction of the structural unit including two layers of Al3B48C2–SiC, one layer of Ti3Si(Al)C2–TiC, and two layers of the in-situ formed TiB2 at the interface. The formation process of Al–B–C phase is also
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displayed. First, Al–Si alloy was infiltrated into the green body by capillary action in a
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vacuum environment. In the TiC layer, TiC particles reacted with Si to form Ti3SiC2; moreover, Al dissolved into Ti3SiC2 to form Ti3Si(Al)C2. In the B4C layer, Al dissolved into large size B4C particles to form the Al–B–C phase due to better wettability of Al than Si. Then, Al–Si alloy would decompose into Al and Si due to
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the consumption of reactions. At the interface of B4C and TiC layers, the B4C and TiC particles reacted to form TiB2 layer, which was due to the lower diffusion rate of Si at
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1300 °C than at 1550 °C thus inhibiting the formation of TiSi2.
Fig. 10. Schematic depiction of in-situ formation of TiB2 and Al–B–C phase of sample S2 infiltrated at 1300 °C with Al40Si60 alloy.
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Fig. 11. ∆Gfo of Eqs. 2 and 6 at different temperatures under the condition of excess Si based on thermodynamic calculations.
Fig. 11 shows the Gibbs free energy of Eqs. 2 and 6 at different temperatures under the condition of excess Si and the values of ∆Gfo increase with increasing temperature. The values of ∆Gfo for Eq. 2 at 1550 °C and Eq. 6 at 1300 °C are −50
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and −310 KJ mol−1, respectively, which tend to drive the Eqs. 2 and 6 in the forward direction. The negative value of ∆Gfo indicates the thermodynamic stability of both
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the reaction products, namely TiSi2 and TiB2, at 1550 and 1300 °C. The highly negative values of ∆Gfo do not indicate that the reactions are more favorable and that
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they are controlled by kinetics at low temperatures [44]. 4. Conclusions
In the present study, the laminated B4C/TiC ceramics with the multilayer
structural unit were fabricated by RMI process with Si and Al40Si60 alloy. The sample S1 infiltrated at 1550 °C with Si exhibited the multilayer structural unit including two layers of B4C–SiC, one layer of Ti3SiC2–TiSi2, and two layers of in-situ synthesized TiSi2 at the interface between B4C and Ti3SiC2–TiSi2 layers. The sample S2 infiltrated 15
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For the sample S1 infiltrated at 1550 °C with Si, the large-sized sample could not be fabricated due to the large volume shrinkage caused by the formation of MAX phase in the TiC layer. The large-sized sample S2 was successfully fabricated by RMI
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process at 1300 °C with Al40Si60 alloy. The thickness of the Al3B48C2–SiC layer,
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TiB2–SiC layer, and TiC–Ti3Si(Al)C2–SiC layer was about 450, 30, and 90 µm, respectively.
The ∆Gfo indicated the thermodynamic stability of the TiSi2 and TiB2. The formation of TiSi2 layer and TiB2 layer between B4C and TiC layers could be
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attributed to the diffusion of elements caused by concentration gradient and liquid phase environment, which promoted the mass transfer. The TiSi2 layer was formed at 1550 °C infiltrated with Si rather than TiB2, which was due to higher diffusion rate of
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Si than B and the concentration gradient effect. The TiB2 layer was formed at the
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interface at 1300 °C infiltrated with Al40Si60 alloy rather than TiSi2 layer, which was due to the low diffusion rate of Si caused by the low temperature and interfaces of Al blocking mass transfer.
The laminated ceramics possessed a low porosity of 2%. The step-like fracture
mode indicated that the crack could be deflected by the multilayer structure with compressive stress layer. The flexural strength and the fracture toughness of sample S2 were 160 ± 2 MPa and 5 ± 0.3 MPa·m1/2, respectively. The Al3B48C2–SiC and
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pressureless sintering in laminated ceramics. Acknowledgments
We would like to thank for the financial support from the Chinese National
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Foundation for Natural Sciences under Contracts (Nos.51632007 and 51672218) and
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the National Key R&D Program of China (No. 2017YFB1103500). We would like to thank the Analytical & Testing Center of Northwestern Polytechnical University for
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SEM test.
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ACCEPTED MANUSCRIPT Captions of figures and tables: Fig. 1. Schematic illustration of the fabrication process of green body. Fig. 2. XRD patterns for laminated B4C/TiC ceramics fabricated by RMI process using Si at 1500 °C. Fig. 3. XRD patterns of laminated B4C/TiC ceramics fabricated by RMI process using Al40Si60
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alloy at 1300 °C. Fig. 4. Optical microscopy images (a) and (b) of sample S1 infiltrated at 1550 °C with Si and (c) and (d) of sample S2 infiltrated with Al40Si60 alloy at 1300 °C. Fig. 5. BSE images of sample S1 infiltrated with Si at 1550 °C Fig. 6 BSE images of sample S2 infiltrated at 1300 °C with Al40Si60 alloy. Fig. 7. (a) EDS mapping results of Al and Si and (b), (c) BSE images of sample S2 infiltrated at
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1300 °C with Al40Si60 alloy. Fig. 8. (a) Micrographs of crack propagation and (b), (c), and (d) fracture surface of sample S2
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infiltrated at 1300 °C with Al40Si60 alloy. Fig. 9. Load-displacement curves of sample S2 infiltrated at 1300 °C with Al40Si60 alloy. Fig. 10. Schematic depiction of in-situ formation of TiB2 and Al–B–C phase of sample S2 infiltrated at 1300 °C with Al40Si60 alloy. Fig. 11. ∆Gfo of Eqs. 2 and 6 at different temperatures under the condition of excess Si based on thermodynamic calculations.
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Table 1 EDS results of the point marked in Fig. 5. Table 2 EDS results of the point marked in Figs. 6 and 7.
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ACCEPTED MANUSCRIPT behaviors of SiBC matrix modified C/SiC composites fabricated by reactive melt infiltration, Carbon, 77 (2014) 886-895. [21] A. Grytsiv, P. Rogl, Aluminium-Boron-Carbon, 11E1 (2009) 10-38. [22] C. Zhang, H. Ru, H. Zong, W. Sun, J. Zhu, W. Wang, X. Yue, Coarsening of boron carbide grains during the infiltration of porous boron carbide preforms by molten silicon, Ceram. Int., 42 (2016) 18681-18691. [23] N. Frage, L. Levin, N. Frumin, M. Gelbstein, M.P. Dariel, Manufacturing B4C-(Al,Si) composite
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[31] C. Wu, Y. Li, S. Xie, Micro-structure, mechanical properties and comparison of monolithic and laminated Ti-B4C composite with Al doped, J. Alloys Compd., 733 (2018) 1-7. [32] H. Foratirad, H. Baharvandi, M.G. Maraghe, Effect of excess silicon content on the formation of nano-layered Ti3SiC2 ceramic via infiltration of TiC preforms, J. Eur. Ceram. Soc., 37 (2017) 451-457. [33] B.Y. Liang, M.Z. Wang, J.F. Sun, X.P. Li, Y.C. Zhao, X. Han, Synthesis of Ti3SiC2 in air using
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[35] L. Wang, X. Yin, X. Fan, P. Greil, N. Travitzky, Ti3Si(Al)C2-based ceramics fabricated by reactive melt infiltration with Al70Si30 alloy, J. Eur. Ceram. Soc., 34 (2014) 1493-1499. [36] C.W. Bale, P. Chartrand, S.A. Degterov, G. Eriksson, K. Hack, R.B. Mahfoud, J. Melançon, A.D. Pelton, S. Petersen, FactSage thermochemical software and databases, Calphad., 33 (2002) 295-311. [37] I. Kero, R. Tegman, M.L. Antti, Effect of the amounts of silicon on the in situ synthesis of Ti3SiC2 based composites made from TiC/Si powder mixtures, Ceram. Int., 36 (2010) 375-379. [38] G. Will, Die Kristallstruktur von C4AlB24, Acta. Crystallogr., 25 (1969) 1219-1222. [39] C. Xiaozhou, W. Chao, X. Xiangxin, C. Gongjin, Effect of Ti Addition on the Residual Aluminium Content and Mechanical Properties of the B4C-Al Composites Produced by Vacuum Infiltration, Arch. Metall. Mater., 60 (2015). [40] Z. Luo, Y. Song, S. Zhang, D.J. Miller, Interfacial Microstructure in a B4C/Al Composite Fabricated by Pressureless Infiltration, Metall. Mater. Trans. A, 43 (2011) 281-293. [41] X. Ma, X. Yin, X. Fan, X. Sun, L. Yang, F. Ye, L. Cheng, Microstructure and properties of dense
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ACCEPTED MANUSCRIPT Tyranno-ZMI SiC/SiC containing Ti3Si(Al)C2 with plastic deformation toughening mechanism, J. Eur. Ceram. Soc., 38 (2017) 1069-1078. [42] M. Sun, Y. Bai, M. Li, S. Fan, L. Cheng, In-situ fabrication of laminated SiC/TiSi2 and SiC/Ti3SiC2 ceramics by liquid silicon infiltration, Ceram. Int., 44 (2018) 11410-11416. [43] R. Bermejo, “Toward seashells under stress”: Bioinspired concepts to design tough layered ceramic composites, J. Eur. Ceram. Soc., 37 (2017) 3823-3839. [44] M. Krinitcyn, Z. Fu, J. Harris, K. Kostikov, G.A. Pribytkov, P. Greil, N. Travitzky, Laminated
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Object Manufacturing of in-situ synthesized MAX-phase composites, Ceram. Int., 43 (2017)
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Highlights: The laminated B4C/TiC ceramics were designed and fabricated by RMI process.
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The Al3B48C2 and AlB24C4 were formed by the reaction between Al and B4C.
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TiSi2 and TiB2 layers were formed at the interface between B4C and TiC layers.
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S0925-8388(18)32392-2
DOI:
10.1016/j.jallcom.2018.06.271
Reference:
JALCOM 46607
To appear in:
Journal of Alloys and Compounds
Received Date: 1 February 2018 Revised Date:
13 June 2018
Accepted Date: 23 June 2018
Please cite this article as: M. Sun, Y. Bai, M. Li, S. Fan, L. Cheng, Structural design of laminated B4C/ TiC composite fabricated by reactive melt infiltration, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.06.271. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. 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.
ACCEPTED MANUSCRIPT Structural design of laminated B4C/TiC composite fabricated by reactive melt infiltration Mengyong Sun, Yuhang Bai, Mingxing Li, Shangwu Fan*, Laifei Cheng*
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Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University, Xi’an, Shaanxi 710072, PR China Abstract:
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Laminated boron carbide/titanium carbide (B4C/TiC) ceramic with multilayer structural unit was fabricated by reactive melt infiltration (RMI) process with silicon
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(Si) and aluminum silicon (Al40Si60) alloy. Laminated B4C/TiC ceramics infiltrated at a low temperature of 1300 °C with Al40Si60 alloy acquired structural unit including two layers of Al3B48C2–SiC, one layer of TiC–Ti3Si(Al)C2, and two layers of in-situ
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synthesized TiB2 at the interface between Al3B48C2 and TiC layers. The thickness of Al3B48C2–SiC layer, TiB2–SiC layer, and TiC–Ti3Si(Al)C2–SiC layer was found to be about 450, 30, and 90 µm, respectively. Reaction products of Al and B4C consisted of
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two types of Al–B–C ternary phases including Al3B48C2 and AlB24C4 attributed to the
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solution of Al in B4C and diffusion of Al into B4C. The diffusion of elements caused by concentration gradient and different diffusion rate of Si led to the formation of TiSi2 layer at 1550 °C and TiB2 layer at 1300 °C between B4C and TiC layers. Keywords: Structural design; Laminated ceramics; In-suit formation; Boron carbide– aluminum
1 *Corresponding author: Laifei Cheng, Ph.D., Professor; E-mail address: [email protected]. Shangwu Fan, Professor; E-mail address: [email protected].
ACCEPTED MANUSCRIPT 1. Introduction Boron carbide (B4C) is high-strength and low-weight ceramic with high hardness in combination with high wear resistance, good impact resistance, high melting point,
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resistance to chemical attack, and good electrical conductivity [1-6]. However, the dense B4C ceramics with low fracture toughness (about 3.7 MPa·m1/2) are usually fabricated by hot pressing (HP) and spark plasma sintering (SPS) at high temperature
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(above 2000 °C) and pressure (30 MPa), which limit their potential application as
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structural materials due to undesirable sample shape and expensive product [7-9]. Therefore, in order to overcome the above mentioned limitations, extensive research efforts have been devoted to the investigation of the pressureless sintering of B4C to fabricate ceramics with large size and complex shape. In pressureless sintering, B4C
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ceramics can also be densified using fine grain powders (<3 µm) and high sintering temperature (above 2150 °C) [10, 11]. In order to enhance the sinterability and mechanical properties of B4C ceramics by pressureless sintering, the addition of
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secondary phases (such as C, B, Si, Al, and Mg) as sintering aids has been carried out
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to decrease the sintering temperature of B4C ceramics [12-15]. Moreover, the reactive melt infiltration (RMI) is a rapid and economical pressureless sintering method to fabricate B4C ceramics by infiltrating molten silicon (silicon or aluminum silicon alloy) spontaneously into porous preforms under vacuum at a low temperature (below 1500 °C) [16-20]. Advantages of RMI include short fabrication time, cost effectiveness, and fabrication of near net shape ceramics. Therefore, understanding the phase equilibrium of the Si or alloy–B4C composites during the permeation
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ACCEPTED MANUSCRIPT process plays an extremely important role in the study of wettability at high temperature and structural design of materials [21]. Zhang et al. [22] fabricated the reaction-bonded B4C ceramics by infiltrating
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molten Si into porous preforms at the temperature of 1450, 1650, and 1750 °C. They found that the B12(B, C, Si)3 phase precipitated at the defective concave surface of the large B4C grains. Frage et al. [23] reported that the Al–Si alloys with Si content higher
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than 40 wt.% could prevent the formation of brittle compound Al4C3. Viala et al. [24,
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25] systematically studied the chemical reactivity of B4C with Al during the RMI process. They found B4C to be unstable and it reacted with either solid or molten Al. When the temperature was above 868 °C, the reaction products were Al3BC and Al3B48C2 (β-AlB12) phases. Lee et al. [26] found numerous reactant phases including
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AlB2, Al3B48C2, AlB24C4 (AlB10), and Al3BC phases at the B4C–Al boundaries by the pressureless infiltration technique at 800 °C for 1 h. Arslan et al. [8] reported that the infiltration process should be performed at higher temperature of 1200 °C due to the
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wettability requirement between Al and B4C.
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Furthermore, layered structure design of ceramics is also an effective way to further improve the fracture toughness [27-31]. The laminated ceramic composites offer the best prospect for rational use of the unique physical and mechanical properties of monolithic ceramics, thus providing a way to improve the durability; toughness; and wear, corrosion, and thermal resistance. Recently, the MAX phase with nano-layered structure has been used to improve the fracture toughness [32-34]. Wang et al. investigated the Ti3Si(Al)C2-based ceramics synthesized by RMI of
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ACCEPTED MANUSCRIPT Al70Si30 alloy into the porous TiC preforms [35]. Therefore, the low temperature preparation by RMI process and multi-scale layer structure design can be combined for the fabrication of the laminated B4C ceramics with high toughness at low cost.
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In this study, the laminated B4C/titanium carbide (B4C/TiC) ceramics were fabricated by tape casting, lamination, and RMI process with silicon (Si) and aluminum silicon (Al40Si60) alloy. The structural design of laminated ceramics was
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investigated to study the interfacial reaction and phase transformation during the RMI
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process. Finally, the microstructures and mechanical properties of the laminated B4C/TiC composite were characterized. 2. Experiments
Tape casting is the most common method to fabricate the laminated ceramics.
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Fig. 1 shows the fabrication process of green body. For tape casting, slurries of compound powders (50 wt.%) were made by mixing powder blends with solvent (mixed alcohol and butanone, 40 wt.%) and dispersing agent (triethylphosphate, 2
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wt.%), and then milled for 24 h. Subsequently, binder (polyvinyl butyral (PVB), 4
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wt.%), plasticizer (dioctyl phthalate (DOP), 2 wt.%), and glycerol (2 wt.%) were added and the slurries were milled for an additional 24 h. Further, the slurries were degassed, and tape casting was carried out using an experimental tape casting equipment (Beijing Orient Sun-Tec Co., Ltd.) with a dual doctor blade casting head. After drying, the B4C–SiC green tape including B4C (75 wt.%, D50 = 20 µm), SiC (20 wt.%, D50 = 20 µm), and graphite powders (5 wt.%, D50 = 1 µm); and the TiC green tape including TiC (D50 = 1 µm) and graphite (10 wt.%) were fabricated. The
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ACCEPTED MANUSCRIPT mass fractions of different powders were based on the quality of the total powders in the green tape. Then, the green tapes were cut into 40 mm × 40 mm pieces with the thickness of 0.2 mm. The laminated green compacts were fabricated by stacking three
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layers of B4C–SiC green tapes and one layer of TiC green tape alternately.
Fig. 1. Schematic illustration of the fabrication process of green body.
Further, the green compact was laminated by low thermal compression at 150 °C
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under a pressure of 30 MPa. Then, the binder removal process was performed at 600 °C for 3 h with a heating rate of 1 °C min−1 in argon flow followed by sintering at
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1300 °C with a holding time of 1 h. Finally, the laminated B4C/TiC compacts with porosity of about 40% were infiltrated under vacuum (10 Pa) at 1550 °C for 1 h with
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Si designated as sample S1 and at 1300°C for 2 h with Al40Si60 alloy designated as sample S2, respectively. Open porosity was measured by the Archimedes’ method. Microstructures of the
samples were systematically characterized by scanning electron microscopy (SEM, S-4700, Japan) and optical microscopy (OM, OLYMPUS, Japan). The elemental composition and spatial distribution of ceramics were characterized by energy dispersive X-ray spectroscopy (EDS). The phase composition of materials was 5
ACCEPTED MANUSCRIPT characterized by the X-ray diffraction (XRD). The bending strength was measured via a three-point bending test using an electromechanical universal testing machine (CMT5504, MTS SYSTEMS, Co., Ltd., China). Fracture toughness (KIC) was
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evaluated by single-edge notched beam (SENB) test. The Vickers hardness was measured by Vickers’ indentation with a 9.8 N load applied for 15 s on polished
the FactSageTM software package (version 5.3) [36].
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3. Results and discussion
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sections (AMH43, LECO, USA). Thermodynamic analysis was carried out by using
The laminated B4C/TiC green bodies were fabricated by RMI process with Si at around 1550 °C and Al40Si60 alloy at 1300 °C. The siliconized sample S1 infiltrated at 1550 °C using Si exhibited large deformation, such as cracks and delamination,
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during the process of cooling, which resulted in fragmentation of the sample. The siliconized samples S2 infiltrated at 1300 °C using Al40Si60 alloy exhibited nearly dense composites with the porosity of 2 %. During the process of infiltration, the
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possible reactions are as follows [32, 37-41]: 3TiC(s) + 2Si(l) → Ti 3SiC2 (s) + SiC(s)
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(1)
TiC(s) + 3Si(l) → TiSi 2 (s) + SiC(s)
(2)
TiSi 2 (s) + Al-Si(l) → Ti-Si-Al(l)
(3)
2TiC + Ti-Si-Al(l) → Ti3Si(Al)C2 (s)
(4)
C(s) + Si(l) → SiC(s)
(5)
B4C(s) + 2 TiC(s) + 3Si(l) → 2TiB2 (s) + 3SiC(s)
(6)
Al(l) + B4 C(s) → Al3 B48C2 (s) / Al B24C4 (s)
(7)
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°C.
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Fig. 2. XRD patterns of laminated B4C/TiC ceramics fabricated by RMI process using Si at 1500
Fig. 3. XRD patterns of laminated B4C/TiC ceramics fabricated by RMI process using Al40Si60 alloy at 1300 °C.
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Figs. 2 and 3 exhibit the XRD patterns of samples S1 and S2, respectively. The
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XRD analysis results revealed that Ti3Si(Al)C2, Ti3SiC2, Al3B48C2, TiB2, and SiC were formed by RMI process. The residual Si, Al, and TiC were also detected. TiSi2 and Ti3SiC2 were formed by the reactions between TiC particles and Si (Eqs. 1 and 2). For the formation of Ti3Si(Al)C2, the ternary Ti–Si–Al melt was formed first, attributed to the dissolution of TiSi2 into Al–Si melt (Eq. 3). The TiC twins were formed due to decreasing boundary energy in the presence of Al, and then the Ti3SiC2 grains nucleated from TiC twins. Further, Al dissolved into Ti3SiC2 to form Ti3Si(Al)C2 (Eq.
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ACCEPTED MANUSCRIPT 4) [41]. The solution and diffusion of Al into B4C resulted in the formation of Al3B48C2 and AlB24C4 (Eq. 7). TiB2 was formed due to the reaction between TiC and
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B4C (Eq. 6).
Fig. 4. Optical microscopy images (a) and (b) of sample S1 infiltrated at 1550 °C with Si and (c) and (d) of sample S2 infiltrated with Al40Si60 alloy at 1300 °C.
Fig. 4 shows the optical micrographs of the laminated B4C/TiC ceramics
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fabricated at different temperatures. The thickness ratio of B4C layer to TiC layer is
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about 3:1. Large size particles of B4C and SiC evenly disperse in the thick layer, and Ti3SiC2 (sample S1) or Ti3Si(Al)C2 (sample S2) locates in the thin layer. Figs. 4(b) and 4(d) show that TiSi2 and TiB2 layers were in-situ formed at the interface between B4C and TiC layers of samples S1 and S2, respectively. Fig. 4(a) exhibits large cracks in the sample S1 generated from the process of cooling, which resulted in breakdown of ceramic into small pieces. As a result, large size of laminated B4C/TiC ceramics
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no cracks are observed in the sample S2 filled with solid Si and Al.
Fig. 5. BSE images of sample S1 infiltrated with Si at 1550 °C
Fig. 5 shows the back-scattered electron (BSE) images of sample S1 and Table 1
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lists the EDS results of the points marked in Fig. 5, clearly indicating the distribution of TiSi2, Ti3SiC2, and SiC phases. In the B4C–SiC layer, the SiC and B4C particles are
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encapsulated by solid Si as shown in Fig. 5(b). Besides, a small amount of B12(C, Si, B)3 was also formed in the liquid Si environment based on EDS result of point 6 as
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shown in Fig. 5(b) [22]. Figs. 5(c) and 5(d) exhibit that the in-situ formed TiSi2 phase is located at the interface between B4C and TiC layers, and also inside the TiC layer. Further, Fig. 5(d) demonstrates almost complete conversion of all TiC phases into Ti3SiC2, TiSi2, and small-sized SiC phases. Formation of TiSi2 layer between B4C layer and TiC layer is attributed to the diffusion of elements caused by concentration gradient, and tensile stress due to the volume shrinkage at the interface [42]. The high temperature of 1550 °C (above the melt point of Si) led to faster diffusion of Si into 9
ACCEPTED MANUSCRIPT the interface than B, which resulted in the formation of TiSi2 (Eq. 2). Table 1 EDS results of the point marked in Fig. 5. 2
3
4
5
6
7
40.21 16.45 43.34
72.67 27.33
45.44 02.17 52.39
53.51 46.49 -
72.05 27.59 0.35 -
45.62 13.83 40.55 -
100.00 -
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B C Si Ti
1
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Point Element (at. %)
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Fig. 6 BSE images of sample S2 infiltrated at 1300 °C with Al40Si60 alloy.
Fig. 6 shows the BSE images of sample S2 infiltrated at 1300 °C with Al40Si60
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alloy. The laminated structure of B4C/TiC ceramic is composed of Al3B48C2–SiC–Si, TiB2–SiC, and TiC–Ti3Si(Al)C2–SiC phases. The thickness of the Al3B48C2–SiC–Si layer, TiB2–SiC layer, and TiC–Ti3Si(Al)C2–SiC layer is about 450, 30, and 90 µm, respectively. The in-situ formed TiB2 layer is located in the interface between B4C and TiC layers, which is attributed to the diffusion of Ti and B into the interface due to concentration gradient in Si–Al melt. Owing to the diffusion and reaction of B4C, the SiC was left on one side of the TiB2 layer. Apparent growth of grains could also be 10
ACCEPTED MANUSCRIPT seen, which is attributed to the enough holding time at 1300 °C for 2 h and the suitable liquid phase environment. Formation of TiB2 layer between TiC layer and B4C layer rather than TiSi2 could be attributed to low diffusion rate of Si caused by
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low temperature of 1300 °C and interfaces of Al blocking mass transfer. Consequently, the concentration gradient of Ti and B made the elements diffuse into the interface, which led to the formation of TiB2 (Eq. 6).
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Fig. 7 shows the EDS mapping results and BSE images of sample S2. Fig. 7(a)
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demonstrates that the residual Al and Si are independent rather than Al–Si alloy, and are mainly located between large size SiC particles. The luminance decreases with decreasing content of Si in Si and SiC phases. Fig. 7(b) shows that in the TiC layer, the phase consists mainly of Ti3Si(Al)C2, a small amount of TiSi2, residual TiC, and
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dendritic-like structure of SiC. In the B4C layer as shown in Fig. 7(c), the phase consists mainly of SiC phase and two types of Al–B–C phases resulting from the Eq. 7. According to the EDS results listed in Table 2, the gray area represents the Al3B48C2
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phase with the high Al content, and the dark gray area denotes the AlB24C4 phase with
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the low Al content. Owing to the heat treatment at 1300 °C and sufficient holding time, the reaction products of Al and B4C include mainly Al3B48C2 phase. Besides, the reaction could be regarded as the solution and diffusion of Al into B4C, thus the AlB24C4 phase was mainly found to be located in the interior of the Al–B–C phase.
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Fig. 7. (a) EDS mapping results of Al and Si and (b), (c) BSE images of sample S2 infiltrated
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at 1300 °C with Al40Si60 alloy. Table 2 EDS results of the point marked in Figs. 6 and 7. Point
1
Element (at.%)
13.18 86.82 -
3
4
48.70 5.99 40.11 5.20
39.66 9.95 10.42 3.73 36.23
47.54 3.72 48.74
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B C Al Si Ti
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5
6
7
30.13 5.39 17.34 47.14
60.11 38.46 1.43 -
73.06 20.29 6.65 -
Fig. 8 shows crack propagation and fracture surface of sample S2. Fig. 8(a)
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demonstrates that the crack is obviously deflected at the interface between the B4C and the TiC layers, which was due to the residual stress caused by different thermal
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expansion coefficients. Figs. 8(b) and 8(c) show that the mode of strong binding occurs at the interface, and the fracture modes include transgranular fracture, and intergranular-fracture. Fig. 8(d) exhibits the distribution of Ti3Si(Al)C2 phase in the TiC layer.
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Fig. 8. (a) Micrographs of crack propagation and (b), (c), and (d) fracture surface of sample S2
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infiltrated at 1300 °C with Al40Si60 alloy.
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Fig. 9. Load-displacement curves of sample S2 infiltrated at 1300 °C with Al40Si60 alloy.
Fig. 9 shows the load-displacement curves with the step-like fracture mode for
sample S2, which could be attributed to the laminated structure with compressive layers. In the laminated ceramics with strong interface, the crack can be deflected by residual compressive stress at the interface, which could improve the toughness. The fracture toughness increases with the increase in the crack length, and this is called R-curve behavior [43]. The flexural strength and the fracture toughness of laminated
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ACCEPTED MANUSCRIPT Al3B48C2–SiC/TiC–Ti3Si(Al)C2 ceramics were 160 ± 2 MPa and 5 ± 0.3 MPa·m1/2, respectively. The Al3B48C2–SiC and TiC–Ti3Si(Al)C2 layers exhibited the Vickers’ hardness of 27.5 ± 0.5 and 15 ± 1 GPa, respectively.
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Fig. 10 exhibits the schematic depiction of the structural unit including two layers of Al3B48C2–SiC, one layer of Ti3Si(Al)C2–TiC, and two layers of the in-situ formed TiB2 at the interface. The formation process of Al–B–C phase is also
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displayed. First, Al–Si alloy was infiltrated into the green body by capillary action in a
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vacuum environment. In the TiC layer, TiC particles reacted with Si to form Ti3SiC2; moreover, Al dissolved into Ti3SiC2 to form Ti3Si(Al)C2. In the B4C layer, Al dissolved into large size B4C particles to form the Al–B–C phase due to better wettability of Al than Si. Then, Al–Si alloy would decompose into Al and Si due to
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the consumption of reactions. At the interface of B4C and TiC layers, the B4C and TiC particles reacted to form TiB2 layer, which was due to the lower diffusion rate of Si at
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1300 °C than at 1550 °C thus inhibiting the formation of TiSi2.
Fig. 10. Schematic depiction of in-situ formation of TiB2 and Al–B–C phase of sample S2 infiltrated at 1300 °C with Al40Si60 alloy.
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Fig. 11. ∆Gfo of Eqs. 2 and 6 at different temperatures under the condition of excess Si based on thermodynamic calculations.
Fig. 11 shows the Gibbs free energy of Eqs. 2 and 6 at different temperatures under the condition of excess Si and the values of ∆Gfo increase with increasing temperature. The values of ∆Gfo for Eq. 2 at 1550 °C and Eq. 6 at 1300 °C are −50
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and −310 KJ mol−1, respectively, which tend to drive the Eqs. 2 and 6 in the forward direction. The negative value of ∆Gfo indicates the thermodynamic stability of both
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the reaction products, namely TiSi2 and TiB2, at 1550 and 1300 °C. The highly negative values of ∆Gfo do not indicate that the reactions are more favorable and that
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they are controlled by kinetics at low temperatures [44]. 4. Conclusions
In the present study, the laminated B4C/TiC ceramics with the multilayer
structural unit were fabricated by RMI process with Si and Al40Si60 alloy. The sample S1 infiltrated at 1550 °C with Si exhibited the multilayer structural unit including two layers of B4C–SiC, one layer of Ti3SiC2–TiSi2, and two layers of in-situ synthesized TiSi2 at the interface between B4C and Ti3SiC2–TiSi2 layers. The sample S2 infiltrated 15
ACCEPTED MANUSCRIPT at 1300 °C with Al40Si60 alloy exhibited the multilayer structural unit including two layers of Al3B48C2–SiC, one layer of TiC–Ti3Si(Al)C2, and two layers of in-situ synthesized TiB2 at the interface between Al3B48C2 and TiC layers.
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For the sample S1 infiltrated at 1550 °C with Si, the large-sized sample could not be fabricated due to the large volume shrinkage caused by the formation of MAX phase in the TiC layer. The large-sized sample S2 was successfully fabricated by RMI
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process at 1300 °C with Al40Si60 alloy. The thickness of the Al3B48C2–SiC layer,
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TiB2–SiC layer, and TiC–Ti3Si(Al)C2–SiC layer was about 450, 30, and 90 µm, respectively.
The ∆Gfo indicated the thermodynamic stability of the TiSi2 and TiB2. The formation of TiSi2 layer and TiB2 layer between B4C and TiC layers could be
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attributed to the diffusion of elements caused by concentration gradient and liquid phase environment, which promoted the mass transfer. The TiSi2 layer was formed at 1550 °C infiltrated with Si rather than TiB2, which was due to higher diffusion rate of
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Si than B and the concentration gradient effect. The TiB2 layer was formed at the
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interface at 1300 °C infiltrated with Al40Si60 alloy rather than TiSi2 layer, which was due to the low diffusion rate of Si caused by the low temperature and interfaces of Al blocking mass transfer.
The laminated ceramics possessed a low porosity of 2%. The step-like fracture
mode indicated that the crack could be deflected by the multilayer structure with compressive stress layer. The flexural strength and the fracture toughness of sample S2 were 160 ± 2 MPa and 5 ± 0.3 MPa·m1/2, respectively. The Al3B48C2–SiC and
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pressureless sintering in laminated ceramics. Acknowledgments
We would like to thank for the financial support from the Chinese National
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Foundation for Natural Sciences under Contracts (Nos.51632007 and 51672218) and
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the National Key R&D Program of China (No. 2017YFB1103500). We would like to thank the Analytical & Testing Center of Northwestern Polytechnical University for
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SEM test.
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ACCEPTED MANUSCRIPT Captions of figures and tables: Fig. 1. Schematic illustration of the fabrication process of green body. Fig. 2. XRD patterns for laminated B4C/TiC ceramics fabricated by RMI process using Si at 1500 °C. Fig. 3. XRD patterns of laminated B4C/TiC ceramics fabricated by RMI process using Al40Si60
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alloy at 1300 °C. Fig. 4. Optical microscopy images (a) and (b) of sample S1 infiltrated at 1550 °C with Si and (c) and (d) of sample S2 infiltrated with Al40Si60 alloy at 1300 °C. Fig. 5. BSE images of sample S1 infiltrated with Si at 1550 °C Fig. 6 BSE images of sample S2 infiltrated at 1300 °C with Al40Si60 alloy. Fig. 7. (a) EDS mapping results of Al and Si and (b), (c) BSE images of sample S2 infiltrated at
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1300 °C with Al40Si60 alloy. Fig. 8. (a) Micrographs of crack propagation and (b), (c), and (d) fracture surface of sample S2
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infiltrated at 1300 °C with Al40Si60 alloy. Fig. 9. Load-displacement curves of sample S2 infiltrated at 1300 °C with Al40Si60 alloy. Fig. 10. Schematic depiction of in-situ formation of TiB2 and Al–B–C phase of sample S2 infiltrated at 1300 °C with Al40Si60 alloy. Fig. 11. ∆Gfo of Eqs. 2 and 6 at different temperatures under the condition of excess Si based on thermodynamic calculations.
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Table 1 EDS results of the point marked in Fig. 5. Table 2 EDS results of the point marked in Figs. 6 and 7.
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Highlights: The laminated B4C/TiC ceramics were designed and fabricated by RMI process.
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The Al3B48C2 and AlB24C4 were formed by the reaction between Al and B4C.
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TiSi2 and TiB2 layers were formed at the interface between B4C and TiC layers.
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