International Journal of Minerals, Metallurgy and Materials Volume 25, Number 4, April 2018, Page 453 https://doi.org/10.1007/s12613-018-1591-x

Evaluation of precipitation hardening in TiC-reinforced Ti2AlNb-based alloys Ya-ran Zhang1), Qi Cai2), Yong-chang Liu1), Zong-qing Ma1), Chong Li1), and Hui-jun Li1) 1) State Key Laboratory of Hydraulic Engineering Simulation and Safety, School of Materials Science and Engineering, Tianjin University, Tianjin 300354, China 2) Materials Genome Institute, Shanghai University, Shanghai 200444, China (Received: 31 August 2017; revised: 29 November 2017; accepted: 8 December 2017)

Abstract: Ti2AlNb-based alloys with 0.0wt%, 0.6wt%, and 2.0wt% carbon nanotube (CNT) addition were fabricated from spherical Ti–22Al–25Nb powder by sintering in the B2 single-phase region. Phase identification and microstructural examination were performed to evaluate the effect of carbon addition on the hardness of the alloys. Carbon was either in a soluble state or in carbide form depending on its concentration. The acicular carbides formed around 1050°C were identified as TiC and facilitated the transformation of α2 + B2 → O. The TiC was located within the acicular O phase. The surrounding O phase was distributed in certain orientations with angles of 65° or 90° O phase particles. The obtained alloy was composed of acicular O, Widmanstatten B2 + O, and acicular TiC. As a result of the precipitation of carbides as well as the O phase, the hardness of the alloy with 2.0wt% CNT addition increased to HV 429 ± 9. Keywords: Ti2AlNb alloy; carbides; microstructure; precipitation hardening; hardness

1. Introduction Titanium and titanium alloys have been widely used in aerospace and automotive manufacturing due to their light weight and excellent high-temperature performance; however, the disadvantages of low hardness, low wear resistance, and high production costs limit their engineering applications to some extent [1–4]. By introducing ceramic-reinforced particles into relatively soft Ti alloys, titanium matrix composites (TMCs) have been developed. These show improved strength, hardness, and high-temperature properties, and have recently become an essential research area for aerospace applications [5–7]. Extrinsic particles generally have the problems of surface contamination, low wettability with the matrix, and their dimensions are difficult to reduce below 1 μm. To address these problems, an in situ particle-reinforced technique with low production costs, clean interfaces, and a strong binding force between particle and matrix is proposed. Among the common reinforced phases, TiC is able to form a stable interface with the matrix and its coefficient of thermal expansion is close to that of the matrix. Combined with its high elasticity modulus, as 4–5 times large as that of Corresponding author: Yong-chang Liu

the matrix, these properties make it an ideal reinforced phase compared with SiC [8–10]. Indeed, it has been proved that a small amount of TiC is enough to significantly improve the hardness and wear resistance of the matrix [11]. Zhang et al. [12] indicated that the morphology of TiC within TMCs was either dendritic or short-rod-like. Such finely dispersed and hard TiC precipitates have been shown to contribute to improvements in hardness, as well as grain refinement in carbon-added Ti153 alloys [13] and various β-Ti alloys [14–15]. Furthermore, Babu et al. [16] attributed the high hardness of Ti–6Al–4V–0.5TiC alloys to the formation of a face-centered-cubic Ti(C,H) metastable phase with high lattice strain. As promising materials for aerospace applications, Ti2AlNb-based alloys benefit from high specific strength, stiffness, and show excellent oxidation and creep resistance [17–19]. Their constituent phases may include the ordered intermetallic orthorhombic (O) phase (Ti2AlNb), the ordered hexagonal-close-packed (hcp) intermetallic α2 phase (Ti3Al), and the body-centered-cubic (bcc) B2/β phase [20–24]. In Ti2AlNb matrix composites, the reinforcement relies on SiC fiber coating or sputtering rather than the dispersion of particles in the matrix, and previous studies have focused

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mainly on interface reactions and/or stability [25–26]. Few studies have evaluated the properties of composites [27] or the effect of TiC on the properties of Ti2AlNb-based alloys. Some attempts have been made to produce Ti2AlNb-based alloys using powder metallurgy. This proves to be an economical fabrication process due to cheap raw materials and low processing costs [28–29]. To date, TMC ingots are obtained by casting, forging, or electric arc melting [30]; those produced from elementary powder via powder metallurgy [31–32] suffer Ti loss as a result of oxidation. Therefore, based on previous trials on the microstructure design of Ti2AlNb-based alloys from pre-alloyed Ti–22Al–25Nb powder [29,33–35], it is postulated that a pre-alloyed Ti2AlNb powder should form TiC-reinforced TMCs via powder metallurgy without oxidation. Hence, we deduced that TiC-reinforced Ti2AlNb matrix composites could be synthesized in situ from pre-alloyed Ti2AlNb powder and CNTs via metallurgy. In the present study, Ti2AlNb alloys with added carbon were prepared using metallurgy from single-B2-phase Ti–22Al–25Nb powder. By comparing the phase content, microstructure, and hardness of CNT-added alloys with and without carbon, the solution effect of carbon and the precipitation effect of the carbides were evaluated. In addition, the types and formation temperatures of the carbides were determined, and the effect of the carbides on phase formation is discussed.

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calculated. To study the effect of CNTs on the phase transformations in Ti2AlNb-based alloys, Ti–22Al–25Nb powder with 50.0wt% CNT addition was heated to 1200°C in a differential scanning calorimetry apparatus (Mettler Toledo TGA/DSC 1) at a heating rate of 20°C·min−1 and the heat flow curve recorded synchronously.

3. Results and discussion Ti2AlNb alloys with 0.0wt%, 0.6wt%, and 2.0wt% CNT addition were sintered at 1200°C for 30 min to investigate the effect of the CNTs on the alloys’ equilibrium phases and structure, the XRD patterns are shown in Fig. 1. The pre-alloyed powder comprised the single B2 phase. In the alloy with 0.0wt% CNTs, after furnace cooling from the single B2 phase region, the equilibrium phase comprised the O phase in addition to the B2 phase. Conversely, the O phase was dominant in the alloys with CNTs. Carbon is an α-phase stabilizing element and an α2 phase appeared in the alloy with 0.6wt% CNT addition. However, it disappeared in the alloy containing 2.0wt% CNTs, and an increase in the amount of the O phase was noted. This suggests the transformation of α2 + B2 → O [26].

2. Experimental Spherical Ti–22Al–25Nb (at%) pre-alloyed powder (approximately 200 µm diameter, single B2 phase) and carbon nanotubes (95.0% purity, 23–30 nm length, 2–5 nm diameter) were mixed in a ratio of Ti–22Al–25Nb + ywt% CNT (y = 0.0, 0.6, and 2) and ball milled for 1 h under an Ar atmosphere. The ball-to-powder ratio was 15:1 and the rotation speed was 360 r/min. Details of the Ti–22Al–25Nb powder can be found in Ref. [26]. After sintering at 1200°C for 30 min, Ti2AlNb alloys with 0.0wt%, 0.6wt%, and 2.0wt% CNTs were obtained by furnace cooling. The phase composition and microstructure of the alloys were characterized using X-ray diffraction (XRD, Bruker D8 Advanced) and scanning electron microscopy (SEM, Hitachi S-4800), respectively. Before the microstructure examination the alloys were chemically etched with Kroll’s reagent. The microhardness measurements were carried out using a Vickers hardness test (1.96 N load) according to the ASTM E18-94 standard. For each of the alloys, eight randomly selected samples were measured and the average hardness value was

Fig. 1. XRD patterns of the sintered Ti2AlNb-based alloys with 0.0wt%, 0.6wt%, and 2.0wt% CNT addition.

Fig. 2 shows the SEM images of the Ti2AlNb alloys with 0.0wt%, 0.6wt%, and 2.0wt% CNT addition. The morphology of each phase is indicated in the images. Different from the near-equiaxial O phase (about 10 μm in length and 5 μm in width) precipitated from the B2 matrix in the alloy without CNTs (Fig. 2(a)), acicular O less than 1 μm in width can be observed in the 0.6wt% CNT-added alloy (Fig. 2(b)). Note that a small amount of α2 was retained and united with the formed O phase. These acicular Os are concentrated in clusters and have an angle of 65° or 90° between them; the orientation relationship between lath Os and the B2 matrix has been reported as [ 111 ]B2//[ 1 11 ]O and (110)B2//(001)O [36].

Y.R. Zhang et al., Evaluation of precipitation hardening in TiC-reinforced Ti2AlNb-based alloys

Furthermore, the white needle-like precipitates, surrounded by this acicular O phase, are assumed to be carbides formed from the pre-alloyed powder and CNTs (Fig. 2(b)). The car-

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bides are about 1–2 μm in size. When 2.0wt% CNTs was added, the amount of acicular O decreased and the amount of precipitate increased (Fig. 2(c)).

Fig. 2. SEM images of the sintered Ti2AlNb-based alloys with 0.0wt% (a), 0.6wt% (b), and 2.0wt% (c) CNT addition.

The carbide and CNT peaks can hardly be discerned in the XRD patterns in Fig. 1. Therefore, a low-rate scanned XRD pattern in certain angle regions was performed on the alloy with 2wt% CNT addition to determine the type of carbide (Fig. 3). Avoiding the high-intensity peaks of Ti2AlNb, it is suggested that the carbides are TiC (Powder Diffraction File #74-1219). The energy dispersive spectrometer (EDS) results (Fig. 3 insets) also indicate that the particles are rich in Ti and C. These particles are located among the α2/O particles, which implies that the generation of TiC is related to the transformation of α2 + B2 → O. The formation of TiC with a large stoichiometric range or carbon solid solution in the α2 phase relies on the diffusion of carbon [37]. The solid solubility of carbon in α-Ti is low, less than 1.8at% at 900°C and even lower than 0.5at% below 600°C [38]. Thus, the carbon atoms were predominantly dissolved in the alloy at 0.6wt% CNTs and when this increased to 2.0wt% the carbon tended to precipitate as carbides, as shown in Figs. 2(b) and 2(c). This stabilization of the carbon could account for the existence of the α2 phase. As the carbon content increased to 2.0wt%, TiC began to precipitate and was surrounded by Nb- and Al-rich areas. The transformation α2 + B2 → O relies on the diffusion of Nb to Nb-depleted regions. The remaining α2 phase and Nb-rich regions transformed into the O phase [39]. Therefore, the precipitation of TiC accelerated the transformation of α2 + B2 → O and α2 was consumed due to the absence of carbon stabilization. In addition, the surrounding B2 matrix was stabilized by the Nb enrichment and the acicular O that transformed from the α2 phase became thinner than that in the CNT-free alloy. Fig. 4 shows the recorded DSC curve for the mixture of Ti–22Al–25Nb powder and CNTs at a heating rate of 20°C·min−1. To reveal the heat flow response of the carbides, 50wt% CNTs were added. Compared with the Ti–22Al–25Nb curve without CNTs [34], an extra exother-

mic peak appears just before the transformation of α2 → B2, i.e., below 1060°C, as indicated by the arrow in Fig. 4. Roger et al. [40] synthesized TiC-reinforced TMCs using powder metallurgy. In their study, complete conversion to TiC was observed at 1200°C in the Ti–CNT mixtures and TiC particle growth was fast. CNTs release free carbon atoms from their top and bottom above 900°C and these tend to combine with Ti forming TiC. This may be the reason for the exothermic peak at 1050°C. Hence, TiC is formed at the heating stage and affects the microstructure of Ti2AlNb-based alloys during cooling.

Fig. 3. XRD pattern of the 2.0wt% CNT-added Ti2AlNb alloy under a low scanning rate. Insets are SEM image of 2.0wt% CNT-added Ti2AlNb alloy and corresponding EDS result.

The volume fraction of the O phase was calculated from the XRD patterns using the Rietveld method [41], Fig. 5. The Vickers hardness is also included in this figure. We can identify an increase in the O phase content and hardness with increasing amounts of CNTs, and O phase precipitation hardening has previously been reported as effective in im-

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proving the hardness of these alloys [42−44]. However, 2.0wt% CNT addition stimulated an enhancement of 100 HV in hardness when the O phase content displayed a limited increment. Fig. 6 shows the high-magnification SEM images of the acicular O in the 0.6wt% and 2.0wt% CNT-added alloys. The average width of the acicular O was 0.42 ± 0.07 μm and 0.41 ± 0.05 μm throughout the samples of the alloys with 0.6wt% and 2.0wt% CNT addition, respectively. In addition, any difference in the Widmanstatten B2 + O structures in the two alloys was hardly identifiable; the widths of the fine O laths were measured as 0.09 ± 0.02 μm and 0.11 ± 0.02 μm, for the 0.6wt% and 2.0wt% CNT-added alloys, respectively. Therefore, we suggest that the improvement in hardness in the 2.0wt% CNT-added alloy, HV 429 ± 9, is related to solution hardening of the carbon and precipitation hardening of the carbides. Considering the transformations during cooling: the α2 phase first precipitated from the B2 matrix, i.e., B2 → α2, followed by the formation of an equiaxial or acicular O phase, i.e., α2 + B2 → O. When the alloys were cooled to the B2 + O phase region, the O phase began to precipitate from the B2 matrix and the Widmanstatten B2 + O structure appeared. Carbon was mostly in a soluble state in the 0.6wt% CNT-added alloy; α2 was retained and united with the formed O phase, due to the α2 stabilization of carbon. The enhancement in hardness is attributed to the refinement of the acicular O phase, the precipitation of fine O laths, and the solution hardening of the carbon. When the content increased to 2.0wt%, the low solubility of carbon in α-Ti resulted in the precipitation of carbides at about 1050°C, i.e., in the α2 phase region. The Ti2AlNb-based alloys experienced the B2, α2 + B2, α2 + B2 + O, and B2 + O phase regions in sequence during cooling [32]. The previously formed carbides served as nucleation sites for the α2 phase at around 1150°C and their presence further facilitated the transformation of α2 + B2 → O. The α2 phase was then completely consumed and the obtained acicular O phase arranged at orientations of 65° or 90° to each

other. As the temperature decreased, a secondary O phase precipitated from the B2 matrix, forming a fine Widmanstatten B2 + O structure. A schematic graph of the microstructure evolution is illustrated in Fig. 7. As the width and content of the O phase in the 2.0wt% CNT-added alloy remained unchanged in contrast with the 0.6wt% CNT-added alloy, the precipitation hardening of carbides is regarded as responsible for the enhancement in hardness.

Fig. 4. DSC curve of the 50.0wt% CNT-added Ti2AlNb alloy with a heating rate of 20°C·min−1.

Fig. 5. Volume fraction of O phase and hardness for the Ti2AlNb-based alloys with 0.0wt%, 0.6wt%, and 2.0wt% CNT addition.

Fig. 6. High-magnification SEM images of the Ti2AlNb-based alloys with 0.6wt% (a) and (b) 2.0wt% CNT addition.

Y.R. Zhang et al., Evaluation of precipitation hardening in TiC-reinforced Ti2AlNb-based alloys

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Fig. 7. Schematic illustration of microstructure evolution for the 2.0wt% CNT-added Ti2AlNb alloy.

4. Conclusion In this study, CNT-added Ti2AlNb alloys were fabricated using powder metallurgy. Carbon was in solution state at 0.6wt% CNT content and carbides precipitated when the CNT concentration exceeded the carbon solubility in α-Ti. The carbides were identified as TiC and were formed just before the transformation of α2 → B2 at the heating stage. An acicular O phase formed around the TiC at angles of 65° or 90° to each other. The solution effect of carbon contributed to the enhanced hardness and was accompanied by precipitation of the O phase in the 0.6wt% CNT-added alloy. The further enhancement in hardness in the 2.0wt% CNT-added alloy is attributed to the precipitation effect of TiC.

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Acknowledgements The authors are grateful to the China National Funds for Distinguished Young Scientists (No. 51325401), the National Natural Science Foundation of China (Nos. 51474156 and U1660201), and the National Magnetic Confinement Fusion Energy Research Program of China (No. 2014GB125006) for financial support.

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