Preparation of Ti3AlC2 bulk ceramic via aqueous gelcasting followed by Al-rich pressureless sintering

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Original Article

Preparation of Ti3AlC2 bulk ceramic via aqueous gelcasting followed by Alrich pressureless sintering Zhanchong Zhaoa, Xianhui Lib, Xian Zengc, Xiaoxin Zhanga, Qingzhi Yana,* a

Laboratory of Special Ceramics and Powder Metallurgy, School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, 100083, China b Luoyang Institute of Science and Technology, No. 90, Wangcheng Avenue, Luolong District, Luoyang City, Henan Province, 471023, China c China Nuclear Power Technology Research Institute, Shenzhen, 518000, China

A R T I C LE I N FO

A B S T R A C T

Keywords: Ti3AlC2 bulk ceramic Aqueous gelcasting Al-rich pressureless sintering C-rich pressureless sintering Mechanical properties

Ti3AlC2 bulk ceramics were prepared via aqueous gelcasting followed by C-rich and Al-rich pressureless sintering. The optimized pH value, zeta potential, dispersant content, and solid loading content were determined to be 10, 72.6 mV, 1.6 wt%, and 52 vol%, respectively. Impurities at ppm level containing in the flowing argon could cause severe decomposition of gelcasted bulk Ti3AlC2, forming whiskers of Al2OC and Al4O4C and floccule of AlN. C-rich pressureless sintering resulted in the delamination of a duplex layer of Ti(CO) and Ti3(AlO)Cx-Ti (O,C). The channels formed after debinding facilitated the outward diffusion of Al and the inward diffusion of O and C, and thereby promoting the decomposition of C-rich sintered Ti3AlC2. The combined effect of the unclosed channels and the porous reaction Ti3(AlO)Cx-Ti(O,C) layer brought a catastrophic reduction in the density and mechanical properties of the C-rich sintered Ti3AlC2 ceramic. While the Al-rich pressureless sintering system isolated C, CO and N2 and supplied a closed Al-rich atmosphere, thereby suppressing the decomposition reactions and promoting the sintering densification and ultimately leading to the superior in mechanical properties. The density, hardness, flexural strength and fracture toughness of the Al-rich sintered ceramic reaches 4.13 g/ cm3, 4.36 GPa, 345 MPa, 4.79 MPa m1/2, respectively.

1. Introduction Ti3AlC2 belongs to the group of layered-ternary compounds called MAX phases or Mn+1AXn, where M is an element of transition metal, A is an element of A-group (mostly from Ⅲ A to Ⅳ A), X is either carbon or nitrogen and n = 1, 2 or 3. And due to its superb compatibility in contact with LBE (lead-bismuth-eutectic) and excellent radiation resistance [1,2], Ti3AlC2 is considered as one of the most promising candidate for high temperature structural materials and is currently being assessed as the fuel cladding and cooling pump impellers in advanced nuclear reactors of accelerator-driven systems (ADS) and Gen-IV lead-cooled fast reactors (Gen-IV LFRs) [3,4]. In the past two decades, dense bulk ceramic of Ti3AlC2 is prepared in the temperature range from 1400 °C to 1600 °C, using pressure-assisted forming technology, such as HP (Hot-Pressing Sintering) [5], RHP (Reactive Hot-Pressing Sintering) [6], SPS (Spark-Plasma Sintering) [7], HIP (Hot-Isostatic Pressing) [8] and SHS/PHIP (Self-propagating High-temperature combustion Synthesis with Pseudo Hot

Isostatic Pressing) [9]. However, these conventional techniques are limited by the simple geometry they yield [10]. And although Ti3AlC2 is a layered machinable ceramic, the machining process leads a huge waste of 90 wt% (an impeller, for example) of the starting cylinder [11]. Considering the synthetic cost of Ti3AlC2 powder, such a level of material waste is unacceptable and a major obstacle to its widespread application. New processing methods need to be explored toward the industrial application of Ti3AlC2. Gelcasting establishes a pressureless and near-net-shape forming method to prepare ceramic components with precise dimensions and complex shapes [10,12]. However, the literature on the preparation of Ti3AlC2 or the MAX phase bulk ceramic via gelcasting and pressureless sintering was very limited. Barsoum et al. [13] reported the tape casting, pressureless sintering and grain growth in the Ti3SiC2 compacts. In this research work the authors found that high C activities, resulting from the graphite heater in the sintering furnace, lead to a loss of Si and the conversion of Ti3SiC2 to TiCx. Hu et al. [14] chose to sinter the gelcasted Ti3AlC2 green body in a tube furnace, thereby avoiding

⁎ Corresponding author at: Xueyuan Road 30, Haidian District, Laboratory of Special Ceramics and Powder Metallurgy, School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, 100083, China. E-mail address: [email protected] (Q. Yan).

https://doi.org/10.1016/j.jeurceramsoc.2020.02.028 Received 30 August 2019; Received in revised form 11 February 2020; Accepted 13 February 2020 0955-2219/ © 2020 Elsevier Ltd. All rights reserved.

Please cite this article as: Zhanchong Zhao, et al., Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2020.02.028

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contamination of the sintering atmosphere by the heater. However, as Hu mentioned in his paper [14], decomposition occurred on the surface of the sintered sample attributed to the surface induced effects. As Chen et al. [15] observed in his work, bulk Ti3AlC2 reacted with a very small amount of oxygen impurity (0.6 ppm) in flowing argon to form Al2CO whiskers in a carbon-containing environment. Barsoum et al. [16] fabricated Mg-Ti2AlC composites by pressureless infiltration of molten Mg into Ti2AlC foam. Zhou et al. [17] carried out the pressureless sintering of bulk Ti3AlC2 ceramic in embedded Al4C3 powder and found that it effectively inhibited the decomposition due to the Al-rich atmosphere Al4C3 provided. Later, inspired by Zhou's work [17], Potoczek et al. [18] performed the Al-rich pressureless sintering of Ti2AlC foams by embedding the samples in a Ti3AlC2 powder bed at 1400 °C. In the present work, inspired by the five lines of research work [13,15–17,19], we extend the idea of sintering protection from the “open molten pool” [13], “pressureless infiltration” [16] and “embedded powder” [17] to a closed environment of Al-rich pressureless sintering system (see later). As far as the author knows, in the sintering temperature range from 1450 °C to 1600 °C, Ti3AlC2 is a highly reactive ternary layered compound that it not only decomposes itself due to the surface induced effects but also reacts with many traces of impurities it comes into contact. As a result, the susceptibility of Ti3AlC2 to thermal decomposition is strongly affected by many complicated factors such as the purities of powders and types of MAX phases, sintering atmosphere, sintering temperature, sintering pressure, purities of inert gases, and the kinds of heating elements used. The challenge for the pressureless sintering of bulk Ti3AlC2 ceramic is therefore to attain high density meanwhile controlling the reaction between Ti3AlC2 and impurities it may contact. And, to date, the decomposition problem of pressureless sintering of bulk Ti3AlC2 has not been solved in the literature of existing research. Our present breakthrough fills this gap and hopes to specify the conditions and process parameters of pressureless sintering of gelcasted green body of Ti3AlC2. Moreover, in view of the similarity of structure and properties of the MAX phase ceramics, the processing parameters and sintering methods of Ti3AlC2 determined by our present work could provide a useful reference for the whole MAX phase family.

Fig. 1. The flowchart of Ti3AlC2 gelcasting process.

controlled rotational velocity of 1000 RPM (revolution per minute) using a polytetrafluoroethylene (PTFE) agitator (with two stirring blades). Secondly, the suspensions were degassed in a vacuum deaeration dish before adding APS (2 wt% of AM), PTZ (0.5 wt% of AM) and CTH (0.5 wt% of AM).Then, the suspensions were casted into a glass tube mold and gelled at room temperature for 1 h.

2.2. Al-rich and C-rich pressureless sintering 2. Experiment Fig. 2 gives the schematic diagram of the Al-rich pressureless sintering system of the gelcasted Ti3AlC2 green body. Prior to heating up, the vacuum of the furnace with a carbon heater was pumped to below 50 Pa and then filled with argon to replace the air in the furnace chamber and the molybdenum cup. To further replace the residual air in the furnace chamber and the molybdenum cup, two additional degassing-filling actions were carried out. Leak hunting was performed to ensure that the furnace chamber and the pipeline was leak-free. Then the temperature was raised to 1550 °C at a heating rate of 5 °C/min, during which the pre-placed aluminum powder melted at 660 °C (melting point) [20], and then above 1000 °C the molybdenum cup was enclosed to form a sealed Al-rich protective atmosphere due to the shrinkage of sintered TiC powder [21]. The cylinder sample sintered at 1550 °C for 3 h by this sintering method is referred to as Al-rich sintered cylinder, denoted as AlRC for brief.C-rich pressureless sintering were performed using the same carbon heater furnace in flowing argon at 1550 °C for 3 h with the same heating rate of 5 °C/min without any protection. The high-purity argon was purchased from Beijing Zhongke Tailong Electronic Technology Co., Ltd. The impurities containing in the argon included O2, N2, and H2O. The impurity details provided by the manufacturer are shown in Table 2. It is believed that they have implemented the national standard GB/T 4842-2017 for the detection of impurities in high-purity argon. Based on the results of other researchers, the saturated vapor pressure of aluminum and carbon at the sintering temperature (1550 °C) are 0.7 torr [22] and 10−8 torr [23], respectively. And the cylinder sample sintered in the C-rich atmosphere is referred to as C-rich sintered cylinder, denoted as CRC for brief.

2.1. Starting materials and gelcasting procedure The starting materials required for the preparation process are shown in Table 1. The flowchart of preparation process of Ti3AlC2 green body is shown in Fig. 1. Ti3AlC2 powder was purchased from Forsman Scientific (Beijing). PAA-NH4, AM, MBAM, APS, PLT, PTZ and CTH were purchased from Shanghai Chemical Reagents, China. For the gelcasting procedure, the premix solutions were firstly prepared by dissolving AM (5 wt%, based on Ti3AlC2 powder), MBAM (10 wt% of AM) and PAA-NH4 (0–2.0 wt%) in deionized water. Then, Ti3AlC2 powder was added to the premix solutions and stirred for 6 h to prepare suspensions with solid loading contents ranging from 46 to 54 vol% at a Table 1 The starting materials required for the preparation process. Starting material

Abbreviated

Purities

Functions

Ti3AlC2 Aluminum Deionized water N,N-Methylenebisacrylamide Polyether defoamer N,N-Dimethylacrylamide Ammonium polyacrylate Ammonium persulfate Phenothiazine Catechol

Ti3AlC2 Al H2O MBAM PLT AM PAA-NH4 APS PTZ CTH

98.5 99.5 — 98.5 99.5 97.5 99.5 99.5 97.5 97.5

Powder Atmosphere Solvent Crosslinker Defoamer Monomer Dispersant Initiator Inhibitor Inhibitor

% % % % % % % % %

2

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Fig. 2. The schematic diagram of the Al-rich pressureless sintering system.

placed in 5 ml centrifugal tube with sealing to prevent the evaporation of the solvent. The pH values of the suspensions were titrated by diluted NaOH and HNO3 solution. After 4 h, sedimentation was observed and recorded. Apparent viscosity of the suspensions with selected solid loading and amount of PAA-NH4 was measured using concentric-cylinder system on a rotational viscometer (NXS-11A, made in Chengdu, China) with shear rate range from 10 s−1 to 700 s−1 at room temperature. Each suspension with a constant pH of 10 was prepared by mechanical stirring for 30 min. All measurements were carried out immediately once each suspension was obtained. The density of Ti3AlC2 sintered components was measured based on the Archimedes principle according to ASTM C373-88 (2006) standard and distilled water was used as medium. Flexural strength was determined by the three-point-bend test with the specimen dimensions of 2 × 3 × 12 mm3 and a span length of 7 mm and a loading speed of 0.5 mm/min. Fracture toughness was determined by the single edge notched beam (SENB) method with the specimen dimensions of 2 × 3 × 12 mm3 and a span length of 7 mm. The hardness was measured using a Vickers hardness tester (MH-6, Taiyasaifu, Beijing).

Table 2 Impurities and their concentrations in argon claimed by the manufacturer. Impurity

O2

N2

H2O

Concentration

0.6 ppm

0.2 ppm

0.8 ppm

2.3. Characterization The size distribution of Ti3AlC2 powders was analyzed by BT-9300H (a laser particle size analyzer) and the morphology was analyzed with a scanning electron microscope (SEM, LEO1450, Ltd., Germany) cooperated with an energy dispersive spectrometer (EDS). The Brunauere Emmete Teller (BET) specific surface area (SSA) measurements was carried out by N2 adsorption-desorption at 77 K using an Autosorb system (QuadraSorb Station 3, Quantachrome). The sample was degassed at 300 °C for 2 h under vacuum condition at 102 Pa. The BETSSA was calculated with the ASiQwin-software using the Brunauer–Emmett–Telle equation in the linear relative pressure range of 0.05–0.25. Ti3AlC2 powder X-ray diffraction (XRD) was carried out by a Rigaku Smart Lab (Rigaku, Rigaku Corporation, Tokyo, Japan) diffractometer using Cu Kα radiation (41 KV and 45 mA) and step scan 0.01°, 10°–90° 2 theta range and a step time of 0.25 s. A Netzsch instrument (STA 449F3) was used in the TG-DSC analysis to determine the debinding temperature. The sample of Ti3AlC2 green body was heated from room temperature to 1000 °C at 10 °C/min in flowing helium. Zeta potential was measured through Zetasizer Nano (Malvern Panalytical Corporation). Ti3AlC2 suspensions (0.1 vol%) were prepared with different weight percentages (based on powder weight) of PAA-NH4 at various pH. The suspensions were stirred for 12 h to break particle agglomerates and reach reaction equilibrium between the Ti3AlC2 particle and PAA-NH4. The pH values of the suspensions were titrated by diluted NaOH and HNO3 solution. Sedimentation tests were performed using 0.1 vol% Ti3AlC2 suspensions. The suspensions were

3. Results and discussion 3.1. Powder characterization Fig. 3 presents the morphology and size distribution of starting Ti3AlC2 particles. The starting Ti3AlC2 presented an irregular shape with BET surface area as 0.764 m2/g, and the layered structure was clearly observed. The presence of a certain number of the fine grains would contribute to an increase in sintering activity and sintered body density. The particle size distribution of starting Ti3AlC2 powder is shown in Fig. 3b. The D50 and D90 of Ti3AlC2 powder was 32.59 μm and 42.30 μm respectively, and a bimodal particle size distribution with peaks at 9.5 and 37.3 μm were observed, which indicates a good gradation of Ti3AlC2 powder and helped to solve the problem about the

Fig. 3. (a) SEM image and (b) particle size distribution of Ti3AlC2 powder. 3

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Fig. 4. Zeta potential of Ti3AlC2 suspension with different amount of PAA-NH4 dispersant.

Fig. 6. Effect of solid loading on rheological properties of Ti3AlC2 suspension with 1.6 wt.% PAA-NH4 and 10 pH.

preparation of slurries with low viscosity and high solid loading [24,25].

suspension be in the range of 9.5–10.5. Fig. 5 shows the sedimentation test of Ti3AlC2 suspension at different pH levels from 2 to 12. As can be seen from Fig. 5, the suspensions in the pH range of 8–11 presented the darkest color. This indicates that the sedimentation of these suspensions is minimal and should be attributed to the higher zeta potential they have. And as expected, the suspensions around its IEP pH of 2–6 exhibited a lighter color. Obviously, the sedimentation results agrees well with the zeta potential characterization in Fig. 4.

3.2. Zeta potential and sedimentation stability Fig. 4 shows the zeta potential curves as a function of pH with and without PAA-NH4 for the Ti3AlC2 suspensions. As can be seen in Fig. 4, without adding any dispersant into the suspension, the IEP (isoelectric point) pH of Ti3AlC2 particle was equal to 5.5, which was a little higher than the 4.7 IEP reported in other research works [26,27]. Nevertheless, the zeta potential curve and IEP of Ti3AlC2 are basically similar to those of other MAX phases including Cr2AlC [28], Nb4AlC3 [29], Ti2AlC [30], Nb2AlC [30], and Ti3SiC2 [31]. The addition of even a small amount of 0.4 wt% PAA-NH4 significantly shifted the IEP from 5.5 to 3.7. The curves presented a wide range of pH (pH > 5) with the presence of dispersant, in which the absolute value of zeta potential exceeded 20 mV and stable dispersion suspensions could be obtained in this pH range according to Vallar’s classification standard [32]. From 0.4 wt% to 1.6 wt%, each 0.4 wt% of PAA-NH4 increased the maximum absolute value of zeta potential. And the maximum zeta potential values for the addition of 0.4 wt%, 0.8 wt%, 1.2 wt% and 1.6 wt% dispersant were 50.6 mV, 56.3 mV, 62.2 mV and 72.6 mV, respectively. And the suspension with zeta potential of 50 mv or more could be identified as good stability according to Vallar [32]. Moreover, all the maximum values were obtained at pH 10. However, the excess PAANH4 of 2.0 wt% may impair the absolute zeta potential. It may be because, according to the Stokes Law and the DLVO theory [33,34], excess electrolyte of PAA-NH4 would compress the diffuse electric double layer due to the increased ionic strength, and thereby decreasing the electrostatic repulsion. Based on the above analysis, the optimal PAA-NH4 content should be determined as 1.6 wt% and the pH value of the

3.3. Rheological properties On the one hand, high solid loading is desirable for the final properties of Ti3AlC2 components, but on the other hand, the high solid loading doesn't facilitate mixing and/or casting in slurry processing due to its inherent low viscosity. Hence it is very crucial to keep the fluidity of slurry while optimizing the solid content. This study was carried out on the suspensions containing 46−54 vol% Ti3AlC2 powder. All of the suspensions were added 1.6 wt% PAA-NH4 and the pH of the suspensions were adjusted to 10. Fig. 6 shows the effect of solid loading on apparent viscosity of Ti3AlC2 stable suspensions after mechanical stirring for 6 h. As can be seen in Fig. 6 that suspensions (46−52 vol% solid loading) presented a shear-thinning behavior and a relative lower viscosity, smaller than 1 Pa·s at 15 s−1, which was suitable for the mold casting. However once the solid loading reached 54 vol%, the viscosity sharply increased. This sharp change in viscosity is a consequence of concentrate system, usually associated with the flocculated state of particles within the liquid [35], which signified that the degree of particle agglomerate of flocs increased in suspensions. And due to the nature of in situ forming process, the microstructure of the Ti3AlC2 casting body is identical with or similar to that of the casting suspensions. Once these flocs appear in the suspensions, they will also exist in

Fig. 5. Sedimentation behavior of Ti3AlC2 aqueous suspension with 1.6 wt% PAA-NH4 at different pH levels from 2 to 12. 4

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Fig. 7. (a) Photographs of Ti3AlC2 green bodies of gelcasted cylinders as prepared; (b) and (c) SEM images of the fracture surfaces of gelcasted cylinders.

the green body. And dense Ti3AlC2 ceramic components and parts with good service performance cannot be prepared by sintering green bodies filled with these flocs. Based on the description above, suspension with optimized values of 10 pH, 1.6 wt% PAA-NH4 and 52 vol% solid loading should be used to cast Ti3AlC2 green body. 3.4. Green body characterization The optimized slurry was poured into a mold of glass tube under vacuum environment. No obvious differential sedimentation occurred in the gelation process completed in a short time of 2–4 min. After the monomers polymerized, the wet green body was demolded and cut into small cylinders. Fig. 7a shows the photograph exhibiting small Ti3AlC2 gelcasted cylinders, as prepared. And Fig. 7b and 7c show the SEM images of the fracture surfaces of those gelcasted cylinders. As seen in Fig. 7a, the fracture surfaces of those cylinders were smooth and tidy, and no defects and cracks were observed. It can be seen from Fig. 7b and c that the large and small grains were stacked together to form a good gradation relationship. And no cracks and agglomerates were observed. As can be clearly observed from Fig. 7c, the crisscross polymer network hinged on each Ti3AlC2 particle [14,31]. And this polymeric network structure provided strength for the green cylinders [14,31]. Benefiting from this bonded network structure, the flexural strength of those cylinders reached 35.59 MPa. The relative density of the green cylinders was 63.17 %, which may be partly due to the bimodal particle size distribution.

Fig. 8. TG and DSC profiles of Ti3AlC2 green body in argon atmosphere.

could be related with the evaporation of physically adsorbed water molecules, which was followed by dehydration and evaporation of chemically adsorbed water molecules. And the second and major weight loss that occurred between 200 °C and 600 °C should be ascribed to the burnout of the cross-linked polymer network, associated with an endothermic peak at 400 °C in the DSC curve. As the temperature increases further, after 600 °C, weight loss does not persist and the debinding is completed. In view of the phase equilibrium effect of residual carbon on the main phase of Ti3AlC2, it is very critical and meaningful to maintain a most thorough debinding process. Moreover, in order to reduce residual carbon, it is necessary to maintain a small amount of organics added. And the weight loss curve of the second stage shows a 6.5 wt% organic addition. According to the TG-DSC result in Fig. 8, a reliable debinding process was established. The temperature was raised to 300 °C at a heating rate of 1.5 °C/min under the vacuum environment of 50 Pa, and the

3.5. TG-DSC analysis and debinding TG and DSC profiles of green body between room temperature and 1100 °C in argon atmosphere with the heating rate of 10 °C/min are shown in Fig. 8. Two distinct weight loss stages were observed in the TG profile. In the first stage from ∼200 °C, the beginning of the weight loss 5

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Fig. 9. Photographs of Ti3AlC2 cylinders: (a) before sintering; (b) as-obtained cylinders, sintered directly in a carbon heater furnace under flowing argon; (c) the cleaned sample of CRC; (d) as-obtained cylinders, sintered in a carbon heater furnace under the protection of Al-rich pressureless sintering system.

temperature was maintained for 2.5 h. Then, the temperature was further increased to 600 °C at a heating rate of 0.5 °C/min, and the temperature was maintained for 12 h. 3.6. Sample sintering, phase evolution, and microstructure Fig. 9 shows the photographs of Ti3AlC2 cylinders before and after sintering at 1550 °C for 3 h. As seen in Fig. 9a, before sintering, the surfaces of the Ti3AlC2 cylinder and the sample tray were neat and tidy without any unknown substances attached to them. However, the surface of the as-obtained sample of CRC was covered with white whiskers and the sample tray was covered with a layer of white flocculent material, as shown in Fig. 9b. After cleaning up (see Fig. 9c), the sample of CRC has obvious delamination, a continuous dense layer of the outer ring and an internal loose cracking layer. In contrast, the as-sintered sample of AlRC has no stratification, floccule and whiskers as shown in Fig. 9d. To determine the phase composition of whiskers and flocculent material and the possible phase evolution of the bulk ceramics, the XRD analysis was performed on the samples in Fig. 9. The results were shown in Fig. 10. As can be seen from Fig. 10a, the starting Ti3AlC2 powder was highly pure. And only the phase of Ti3AlC2 (JCPDS#52-0875) was detected on the surface and interior of the AlRC sample according to Fig. 10b and c, which demonstrated that no phase evolution occurred during the whole Al-rich pressureless sintering process. In contrast, two phases of Ti3AlC2 JCPDS#52-0875 and TiC ICSD#44494 were detected from the interior of the CRC sample and three aluminum-containing crystal phases, Al2OC (JCPDS#36-0148), Al4O4C ICSD#18204 and AlN ICSD#44107, were observed in the white whiskers and flocculent material as shown in Fig. 10d and e, respectively. Which suggested that some kind of phase evolution took place during the C-rich pressureless sintering process. Based on the XRD characterization results, the possible reactions involved during the C-rich pressureless sintering process were assumed to be as follows, Ti3AlC2(s) → 2TiC(s) + Al(g) + Ti(g)

(1)

2C(s) + O2(g) → 2CO(g)

(2)

2Al(g) + CO(g) → Al2OC(s)

(3)

Fig. 10. XRD patterns of starting Ti3AlC2 powder and sintered cylinders: (a) starting powder; (b) surface of AlRC; (c) interior of AlRC; (d) interior of CRC; (e) white whiskers and flocculent material collected from surface of CRC sample.

8Al(g) + 4CO(g) → Al4O4C(s) + Al4C3(s)

(4)

Al4C3(s) → 4Al(g) + 3C(s)

(5)

2Al(g) + N2(g) → 2AlN(s)

(6)

Ti3AlC2(s) → 2TiC(s) + Al(g) + Ti(s)

(7)

Mn+1AXn(s) → nMX + A(g) + M(s).

(8)

Based on Low’s [36] systematic researches on the thermal dissociation of bulk Ti3AlC2 in vacuum, the forming mechanism or pathway of TiC formed in the interior of CRC sample could be depicted by Eq. (1). And the forming mechanism or pathway of Al2OC and Al4O4C could be expressed by Eq. (2), Eq. (3) and Eq. (4) according to Gai's research results [15,37]. However, different from Chen's study, Al4C3 was not found in the present study, but AlN was detected by the 6

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XRD analysis. The reason may be that Al4C3 decomposed at a sintering temperature higher than 1400 °C and the formation of AlN could be attributed to the direct combination of N2 and gaseous Al, which could be described by Eq. (5) and Eq. (6), respectively. Furthermore, the formation and growth mechanism of Al2OC and Al4O4C whiskers could be explained by a vapor-liquid-solid (VLS) process [38], which was mainly composed of the reactions between CO and gaseous Al escaped from Ti3AlC2 matrix. Firstly, the gaseous Al escaped outwardly from the matrix through Eq. (1). Secondly, Al gas aggregated and condensed into Al vapor and then turned into tiny aluminum droplets. Thirdly, CO dissolved itself into the tiny aluminum droplets and reacted with aluminum to synthesize whiskers of Al2OC and Al4O4C. In this process, the tiny aluminum droplets, as the initial nucleation site, played a key role in starting and promoting the whiskers growth. In terms of morphology, the white flocculent material formed by sticking to the sample tray should be AlN. And the formation of AlN obeyed the general law of chemical vapor deposition and the sample tray served as the preferential nuclei site for the vapor deposition of AlN. Thus, both CO and N2 could accelerate the thermal dissociation of bulk Ti3AlC2. In addition, titanium nitride (TiN) with a distinctive golden color was not detected on the surfaces of the CRC sample and the sample tray, which made us suspect that only Al escaped from the Ti3AlC2 matrix during the thermal dissociation and Ti did not evaporate through Eq. (1), that is, Eq. (1) should be corrected to Eq. (7). This is inconsistent with Low's study on the thermal dissociation of bulk Ti3AlC2 in vacuum. The discrepancy maybe due to that the gas-phase products of Al and Ti had the similar kinetic characteristics of mass transfer in vacuum, while in argon, the mass transfer efficiency of Al was much better than Ti. Obviously, the melting point of aluminum (660 °C) [39] was much lower than that of titanium (1660 °C). Therefore, at the sintering temperature of 1550 °C, the gas-phase product was mainly Al, not Ti. In general, the surface effects of the MAX phases at the sintering temperature in argon, through the direct evaporation of M and A along with the concomitant formation of the binary carbide (MX), can be expressed as Eq. (8). And based on roughly the same synthetic path and mechanism, the melting point of M is usually much higher than that of A [40]. The gas-phase products are mainly A rather than M. This helps us reasonably make further assumption that the A-rich sintering can suppress surface decomposition during the pressureless sintering of all

Table 3 EDS analysis result of the regions indicated by the yellow arrows in Fig. 11. Element

C (at.%)

Al (at.%)

O (at.%)

Ti (at.%)

Region Region Region Region Region Region Region

18.64 20.57 18.82 22.15 33.08 29.76 31.14

11.66 10.12 – – 17.60 16.42 17.49

10.39 10.94 6.02 10.20 – – –

59.32 58.37 75.15 67.66 49.32 53.82 51.37

1 2 3 4 5 6 7

MAX phases in argon atmosphere. Fig. 11. shows the SEM images of the fracture surfaces of CRC and AlRC. As seen in Fig. 11a and b, the distinctive porous network structure with channel diameters ranging from 5 μm to 20 μm was not completely closed and uniformly distributed in the matrix of CRC sample. The EDS analysis of Region 1 and Region 2, far away from the porous network structure, show that their compositions of chemical elements are Ti, Al, O and C and the atomic ratio is close to 6:1:1:2 as shown in Table 3. While the EDS results of Region 3 and Region 4 near this porous network structure indicate that their compositions of chemical elements are Ti, O and C and the atomic ratio is close to 7:1:2. Combining the XRD analysis from Fig. 10d, Region 1 and Region 2 should be Ti3AlC2, and Region 3 and Region 4 should be TiC. Moreover, in view of the presence of a significant amount of O element, the detected phases of Ti3AlC2 and TiC in the interior of the CRC sample maybe have the structure of Ti3(AlO)Cx and Ti(OC) respectively, which resulted from the reactions between Ti3AlxC2/TiCx/Ti and CO/O [37,41]. The difference of the lattice parameters between TiC and Ti (OC) are very small [41], and so too are Ti3(AlO)Cx and Ti3AlC2 [42]. Consequently, their XRD patterns are similarly presented. Nevertheless, under the premise of Al escaping, a total weight gain of ∼4.5 % was observed after C-rich sintering, which may serve as a compelling evidence for the inward diffusion of O and C from atmosphere to the CRC matrix. In addition, it has been reported that the thermal dissociation was determined by the highly restricted outward evaporation of Al from the matrix to the surface of the bulk MAX phase [43]. Since the kinetic conditions of Al diffusion on the surface and inside of the sample were

Fig. 11. SEM images of the fracture surfaces of the sintered Ti3AlC2 cylinders: (a) and (b) fracture surface of CRC sample; (c) and (d) fracture surface of AlRC sample. 7

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the two samples. C-rich sintering resulted in the uniformly distributed porous network structure with unclosed channels in the interior of the CRC samples. Taking advantage of the convenience provided by these channels, the outward diffusion of Al and the inward diffusion of C and O formed a porous mixture layer of Ti(CO)-Ti3(AlO)Cx. And the combined effect of these factors brought a catastrophic reduction in the density and mechanical properties of the CRC sample. While the Al-rich pressureless sintering system isolated C, CO and N2 and supplied a closed Al-rich atmosphere, promoting the sintering densification and ultimately leading to the superior in mechanical properties.

Table 4 Mechanical properties of Ti3AlC2 cylinders sintered in different atmosphere. Ti3AlC2 cylinder 3

Density (g/cm ) Hardness (GPa) Flexural strength (MPa) Fracture toughness (MPa·m1/2)

AlRC

CRC

4.13 4.36 345 4.79

3.67 4.88 157 3.29

different, the surface Al diffused outwardly more easily than the Al inside the sample. And the density of Ti3AlC2 is smaller than that of TiC, so the phase evolution from Ti3AlC2 to TiCx could lead to the volume shrinkage, which eventually resulted in the delamination of a duplex layer as shown in Fig. 9c. Considering that the outer-ring layer was more conducive to the Al outward diffusion, its phase composition was dominated by Ti(CO). And the cracking layer was a mixture of porous Ti3(AlO)Cx and Ti(O,C), which concurs well with the XRD, SEM and EDS results in this present study. Besides, it is also worth noting that both the orderly arranged pores parallel to the (0 0 0 2) planes reported by Chen et al. [37] and the randomly distributed pores reported by Low et al. [36] cannot be observed in this work. In contrast, the distinctive porous network structure formed after debinding was uniformly distributed in the interior of the CRC sample. The discrepancy between the findings of the present study and the works of Chen et al. and Low et al. may be attributed to the fact that the channels in the porous network structure have already provided the necessary convenience for the outward diffusion of Al and the inward diffusion of O and C in the flowing argon atmosphere, and thereby promoting the decomposition of the CRC sample. This hypothesis got further support from the observation of the fracture surface of the AlRC sample. As can be seen from Fig. 11c and d, the channels in porous network structure were completely closed and the grain boundaries were clearly visible. The crystal grains were closely connected with each other by the boundaries, and the average grain size grew up to about 20.7 μm. The distinctive layered structure in the fracture morphology was clearly visible at high magnification. Besides, the fracture mechanism of the AlRC sample was a mixture of transcrystalline and intergranular type. EDS results of Region 5, Region 6 and Region 7 indicate that their compositions of chemical elements are Ti, Al and C and the atomic ratio is close to 3:1:2. Combining with the XRD analysis from Fig. 10c, those crystal grains should be Ti3AlC2. This analysis results indicate that the Al-rich pressureless sintering system could provide a strong and effective protection for the pressureless sintering of the gelcasted Ti3AlC2 green body. In fact, for the pressureless sintering of Ti3AlC2, it is only necessary to supplement the Al pressure of 0.7 torr (saturated vapor pressure at 1550 °C) in atmosphere to suppress the outward evaporation of Al from bulk Ti3AlC2, thereby suppressing the decomposition reactions of Eq. (1) and/or Eq. (7). Furthermore, in the closed Al-rich pressureless sintering system, the sintering atmosphere can be well controlled and regulated, not only can supplement Al, but also can isolate C, O2, N2 and other impurities, thereby inhibiting the reactions of Eqs. (2)–(6).

4. Conclusion Ti3AlC2 bulk ceramics were prepared by aqueous gelcasting followed by C-rich and Al-rich pressureless sintering. The processing parameters were optimized to obtain stable slurries with low viscosity and high solid loading. The microstructure, sintering mechanism, and mechanical properties of the AlRC and CRC samples were compared and analysed. The conclusions are as follows: (1) The optimized processing parameters of Ti3AlC2 slurries were 10 pH, 1.6 wt% PAA-NH4 and 52 vol% solid loading. The IEP of Ti3AlC2 particle was observed at pH of 5.5, and the highest zeta potential reached 72.6 mV. The sedimentation of Ti3AlC2 suspension is minimal in the pH range 8–11. (2) A small amount of impurities containing in the flowing argon could promote the thermal dissociation of the gelcasted bulk Ti3AlC2, forming whiskers of Al2OC and Al4O4C and floccule of AlN and simultaneously resulting in the delamination of a duplex layer of Ti (CO) and Ti3(AlO)Cx-Ti(O,C). (3) The unclosed channels formed after debinding facilitated the outward diffusion of Al and the inward diffusion of O and C, and thereby promoting the thermal decomposition of the C-rich sintered Ti3AlC2. While the Al-rich pressureless sintering system can provide a strong and effective protection for the gelcasted Ti3AlC2 green body. (4) The combined effect of the unclosed channels and the porous reaction Ti3(AlO)Cx-Ti(O,C) layer brought a catastrophic reduction in the density and mechanical properties of the C-rich sintered Ti3AlC2 ceramic. The density, hardness, flexural strength and fracture toughness of the Al-rich sintered Ti3AlC2 ceramic reached 4.13 g/ cm3, 4.36 GPa, 345 MPa, 4.79 MPa m1/2, respectively”. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. References [1] T. Lapauw, B. Tunca, J. Joris, A. Jianu, R. Fetzer, A. Weisenburger, J. Vleugels, K. Lambrinou, Interaction of M(n+1)AX(n) phases with oxygen-poor, static and fast-flowing liquid lead-bismuth eutectic, J. Nucl. Mater. 520 (2019) 258–272. [2] J. Wang, S. Liu, D. Ren, T. Shao, P. Eklund, R. Huang, Y. Zhu, F. Huang, S. Du, Z. Wang, J. Xue, Y. Wang, Q. Huang, Microstructural evolution of epitaxial Ti3AlC2 film on sapphire under ion irradiation and nanoindentation-induced deformation, J. Nucl. Mater. 509 (2018) 181–187. [3] J.G. Gigax, M. Kennas, H. Kim, T. Wang, B.R. Maier, H. Yeom, G.O. Johnson, K. Sridharan, L. Shao, Radiation response of Ti2AlC MAX phase coated Zircaloy-4 for accident tolerant fuel cladding, J. Nucl. Mater. 523 (2019) 26–32. [4] K. Lambrinou, T. Lapauw, A. Jianu, A. Weisenburger, J. Ejenstam, Corrosion-resistant ternary carbides for use in heavy liquid metal coolants, A Collection of Papers Presented at the 39th International Conference on Advanced Ceramics and Composites, John Wiley & Sons, Inc., 2015, pp. 19–34. [5] J.-H. Han, S.-S. Hwang, D. Lee, S.-W. Park, Synthesis and mechanical properties of Ti3AlC2 by hot pressing TiCx/Al powder mixture, J. Eur. Ceram. Soc. 28 (5) (2008) 979–988. [6] Y. Du, J.-X. Liu, Y. Gu, X.-G. Wang, F. Xu, G.-J. Zhang, Anisotropic corrosion of Ti2AlC and Ti3AlC2 in supercritical water at 500 ºC, Ceram. Int. 43 (9) (2017)

3.7. Density and mechanical properties The density and mechanical properties of the Ti3AlC2 samples of AlRC and CRC were measured and compared. The result was shown in Table 4. As can be seen from Table 4, except hardness, the density, flexural strength and fracture toughness of AlRC are higher than those of CRC. The density, hardness, flexural strength and fracture toughness of AlRC sample reaches 4.13 g/cm3, 4.36 GPa, 345 MPa, 4.79 MPa m1/2, respectively. The reason for the higher hardness of CRC may be that the decomposition product of TiC has a higher hardness than Ti3AlC2. Compared with CRC, the superiority of mechanical properties of AlRC should be attributed to the difference of sintering atmosphere between 8

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