nanosized TiB2–TiC particles

Materials Science & Engineering A 764 (2019) 138266

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A new approach for improving the elevated-temperature strength and ductility of Al–Cu–Mg–Si alloys with minor amounts of dual-phased submicron/nanosized TiB2–TiC particles

T

Yu-Yang Gaoa,b, Feng Qiua,b,∗, Qing-Long Zhaoa,b, Qi-Chuan Jianga,b,∗∗ a

State Key Laboratory of Automotive Simulation and Control, Jilin University, PR China Key Laboratory of Automobile Materials, Ministry of Education and Department of Materials Science and Engineering, Jilin University, Renmin Street NO. 5988, Changchun, Jilin Province, 130025, PR China

b

A R T I C LE I N FO

A B S T R A C T

Keywords: Dual-phased and submicron/nanosized particles θ′ and Q′ precipitates Strength and ductility

At high temperatures, the weakened strength and decreased ductility of Al alloys limit their industrial applications. (TiB2–TiC)/Al–Cu–Mg–Si composites were synthesized by adding in situ dual-phased and bimodal-sized (TiB2–TiC)/Al master alloys to molten Al–Cu–Mg–Si alloys. The addition of minor amounts of TiB2 and TiC particles (0.05 and 0.1 wt%, respectively) effectively refined the α-Al grains and θ′ and Q′ precipitates and increased the elevated-temperature strength and uniform elongation of the Al–Cu–Mg–Si alloys. The 0.1 wt% (TiB2–TiC)/Al–Cu–Mg–Si composite showed the best yield strength (279 MPa), ultimate tensile strength (366 MPa), and uniform elongation (10.6%) at 493 K, with enhancements of 9.4%, 15.1%, and 24.7% compared to the Al–Cu–Mg–Si matrix alloy (255 MPa, 318 MPa, and 8.5%). The simultaneous increases in the elevated temperature strength and ductility of the composites were attributed to the strengthening effects of the bimodalsized TiB2 and TiC particles, θ′ and Q′ precipitation strengthening, and the refined partial recrystallization microstructure.

1. Introduction The weakening of the strength and the decrease in the ductility of Al alloys at elevated temperatures limit their industrial applications [1–4]. The addition of nanosized ceramic particles increased the tensile strength and ductility of Al alloys at elevated temperatures [5–9]. Han et al. prepared in situ TiB2/Al–12Si composites by a mixed-salt reaction method [10]. At 473 K, the yield strength (189 MPa) and ultimate tensile strength (233 MPa) of a 4 wt% TiB2/Al–Si composite increased by 2.7% and 3.6%, respectively, compared with a matrix alloy, while the fracture strain was reduced by 25% [10]. However, nanosized particles tend to form agglomerations, which reduces the enhancement effects, especially in the composites prepared by casting methods. As reported, the addition of a small amount (< 1 wt%) of nanosized particles effectively refined the α-Al grains and increased the strength and ductility of Al alloys at elevated temperatures [11–13]. Tian et al. reported that at 493 K, the yield strength, ultimate tensile strength, and fracture strain of a 0.7 wt% nano-TiC/ZL205A composite were 14.8%,

19.6%, and 84% higher than those of a ZL205A alloy, respectively [13]. Determining how to use small amounts of nanosized particles to greatly increase the strength and ductility of Al alloys at high temperatures has stimulated further studies. Al matrix composites reinforced with bimodal-sized (micron/submicron + nano) ceramic particles achieve more uniformly distributed particles, better thermal physical properties, and better mechanical properties than those reinforced with single-sized particles [5,7]. The bimodal-sized (15 μm + 40 nm) SiC/Al2014 composites had a more homogeneous distribution of nano- and micron-sized particles compared to the single-sized (15 μm/40 nm) SiC/Al2014 composites [7]. Tian et al. reported that the creep resistance of a bimodal-sized (1.88 μm + 97 nm) TiCp/Al–Cu composite was 3–6 times higher than that of the single-sized (1.88 μm/97 nm) TiCp/Al–Cu composites at 453–493 K [5]. On the other hand, the synergistic effects between the different kinds of particles are also beneficial for improving the distribution of the particles and increasing the tensile properties of the Al matrix composites [14,15]. Hu et al. reported that a dual-phased

∗ Corresponding author. Key Laboratory of Automobile Materials, Ministry of Education and Department of Materials Science and Engineering, Jilin University, Renmin Street NO. 5988, Changchun, Jilin Province, 130025, PR China. ∗∗ Corresponding author. Key Laboratory of Automobile Materials, Ministry of Education and Department of Materials Science and Engineering, Jilin University, Renmin Street NO. 5988, Changchun, Jilin Province, 130025, PR China. E-mail addresses: [email protected] (F. Qiu), [email protected] (Q.-C. Jiang).

https://doi.org/10.1016/j.msea.2019.138266 Received 21 May 2019; Received in revised form 6 August 2019; Accepted 8 August 2019 Available online 09 August 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.

Materials Science & Engineering A 764 (2019) 138266

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Fig. 1. (a) The FESEM image, (b) XRD pattern, (c) TEM micrograph and (d) selected-area electron diffraction (SAED) pattern of the in situ (TiB2–TiC)/Al master alloy and (e) the size distribution map of the TiB2 and TiC particles.

involved solution treatment (778 K/2 h), which was followed immediately by quenching in water and aging (433 K/18 h). The phase constitutions of the in situ (TiB2–TiC)/Al master alloy were identified via X-ray diffraction (XRD, D/Max 2500PC Rigaku, Japan). The microstructures were observed by SEM (Tescan vega3 XM, Czech Republic) coupled with EBSD (Oxford NordlysMax) and energy dispersive spectroscopy (EDS), field emission scanning electron microscope (FESEM, JSM 6700 F, Japan), and TEM (JEM-2100 F, Japan). The specimens prepared for the EBSD analysis were first mechanically polished and then electrolytically polished in a solution of 10% HClO4 90% ethanol at 243 K. Step sizes of 2 μm and 0.8 μm were used to characterize the microstructure of the samples in the as-cast and heattreated states, respectively. The EBSD data were analyzed using HKL CHANNEL 5 software. In the EBSD maps, the low-angle grain boundaries (LAGBs) with misorientation angles between 2° and 15° were marked by gray lines, while the high-angle grain boundaries (HAGBs) with misorientation angles greater than 15° were marked by black lines. The spurious grain boundaries with misorientation angles less than the critical value (2°) were not studied. The prepared 0.05 wt% and 0.1 wt% (TiB2–TiC)/Al–Cu–Mg–Si composites were denoted as BC-0.05 and BC-0.1, respectively. The sheets were processed along the extrusion direction to prepare flat dog bone tensile specimens with a 4 × 2 mm2 cross-sectional area and a 10 mm gauge length. The tensile tests were performed by an MTS-810 material testing system (298 K) and Instron-5869 material testing machine (453 K and 493 K) at a constant strain rate of 3 × 10-4 s-1.

(B4C + Al3Ti)/Al composite had a lower level of porosity, more uniform distribution of particles, and improved tensile strength and ductility compared to a single-phased B4C/Al composite [15]. By using the dual-phased and bimodal-sized particles to reinforce the composite, a new generation of Al matrix composites with excellent tensile properties at room and elevated temperatures are expected to be developed. However, the effect of adding minor amounts of bimodal-sized (submicron + nano) dual-phased TiB2 and TiC particles on the microstructure and mechanical properties of Al matrix composites prepared by casting methods has not been reported. In this study, in situ dual-phased (TiB2–TiC)/Al master alloys with particle sizes of 30–730 nm were used. The Al–Cu–Mg–Si composites were prepared by stir casting and master alloys assisted with ultrasonic vibration. The effect of adding minor amounts (0.05/0.1 wt%) of TiB2 and TiC particles on the microstructure, elevated-temperature tensile properties, and the strengthening mechanisms of the composites were discussed. This study provides a new and inexpensive method for the industrial production of Al alloys with excellent elevated-temperature tensile strength and ductility. 2. Experimental procedures The composition of the Al–Cu–Mg–Si alloy was 5.0 Cu, 0.6 Mg, 0.5 Si (wt.%) balanced with Al. The 70Al-21.7Ti-8.3B4C powder system (nominal composition in wt.%) was used to prepare the 30 wt% (TiB2–TiC)/Al alloy. The B4C powder was pretreated by high-speed ball milling at a ball milling speed of 250 rpm, a ball milling time of 20 min, and a ball-to-powder ratio of 50:1. The details of the method used to prepare the in situ dual-phased and bimodal-sized (TiB2–TiC)/Al master alloy have been reported in Ref. [6]. The molar ratio of TiB2 to TiC was 2:1. The weight percentage of TiB2 and TiC particles in the prepared master alloy was 21% and 9%, respectively. When the molten Al alloy was heated to 1123 K, the master alloy was added, which was followed by 5 min of ultrasonication with a vibration power of 5 kW and an ultrasound frequency of 20 kHz. The homogenized rod (φ82 mm × 100 mm) was extruded at 723 K to form a sheet with a cross-sectional area of 8 × 50 mm2. The T6 heat treatment of the sheets

3. Results and discussion Fig. 1(a) and (b) shows the FESEM image and XRD results of the (TiB2–TiC)/Al master alloy, respectively. The master alloy was mainly composed of α-Al, CuAl2, TiB2, and TiC phases without an intermediate Al3Ti phase. Fig. 1(c) and (d) shows the TEM micrograph and corresponding SAED pattern of the master alloy. The distribution of the TiB2 and TiC particles is relatively uniform in the master alloy. Fig. 1(e) shows the statistical diagram of the particle distribution. The number percentages of the nanosized particles (0–100 nm) and submicron-sized 2

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Fig. 2. The as-cast microstructure EBSD maps and SEM images of (a and d) the Al–Cu–Mg–Si alloy, (b and e) the BC-0.05 composite, and (c and f) the BC-0.1 composite. (g) and (h) are the EDS results obtained for point A and B in (f), respectively.

increased the σ0.2, σUTS, and UE of the composites at elevated temperatures. The σ0.2, σUTS, and UE of the BC-0.1 composite were 17.9%, 15.7%, and 5.5% higher than those of the Al alloy at 453 K, respectively, and were 9.4%, 15.1%, and 24.7% higher than those of the Al alloy at 493 K. As indicated in Fig. 4(d), the BC-0.1 composite exhibited a higher work hardening capacity than the Al alloy, which might be because the fraction of fine recrystallized grains in the composite (46.0%) was higher than that of the Al alloy (29.1%). The increase in the work hardening rate indicates that the work hardening capacity of the dislocation accumulations increases, which helped to increase the uniform elongation and delay necking [16,17]. As shown in Fig. 5(a), the BC-0.1 composite with a small amount (0.1 wt%) of TiB2–TiC particles had a relatively ideal strength at 493 K compared to the strength reported for Al–Cu matrix composites [5,6,8,9,13,18–21]. Although the particle content in the BC-0.1 composite was much lower than that in the 40 wt% (TiC–TiB2)/2014Al composite (40BCE) [18], the UE of the BC-0.1 composite was 282% higher than that of the 40BCE composite, while the σUTS of the BC-0.1 composite was only 15 MPa lower than that of the 40BCE composite, as observed in Fig. 5(a) and (b). According to the reported tensile properties of Al–Cu matrix composites at elevated temperatures, the BC-0.1 composite had the best comprehensive tensile strength and uniform elongation at 493 K, as observed in Fig. 5(c). Furthermore, the BC-0.1 composite had the highest σUTS and UE compared to the Al–Cu composites prepared by casting methods. Fig. 6 and Fig. 7 show the TEM micrographs and size distribution maps of the θ′ and Q′ precipitates in the Al alloy and BC-0.1 composite

particles (> 100 nm) were 87.2% and 12.8%, respectively. Fig. 2(a–c) and (d-f) shows the as-cast microstructure EBSD maps and SEM images of the Al–Cu–Mg–Si alloy and (TiB2–TiC)/Al–Cu–Mg–Si composites, respectively. The sizes of the α-Al grains were refined from 4.5 to 402 μm in the Al alloy and 3.4–214 μm in the BC0.05 composite to 3.3–148 μm in the BC-0.1 composite. As indicated in Fig. 2(g) and (h), it can be inferred that the white phase (point A) and the light gray phase (point B) corresponded to the Al2Cu phase and AlCuMgSi phase, respectively. Fig. 3(a–c) shows the EBSD maps of the extruded Al alloy and the composites after the T6 heat treatment. The unrecrystallized coarse grains (50–470 μm) elongated in the extrusion direction. The fine grains (2.6–26.7 μm) formed by dynamic recrystallization in the hot extrusion process were distributed along the grain boundaries (GBs) of the coarse grains. The sizes of the α-Al grains decreased from 3.8 to 345 μm in the Al alloy and 3.1–307 μm in the BC-0.05 composite to 2.6–322 μm in the BC-0.1 composite. As shown in Fig. 3(d–f), the fractions of small grains (2.6–5.0 μm) in the Al alloy, BC-0.05 composite, and BC-0.1 composite were 29.1%, 44.9%, and 46.0%, respectively. Fig. 4(a–c) shows the engineering stress-strain curves of the Al alloy and composites. The yield strength (σ0.2), ultimate tensile strength (σUTS), and uniform elongation (UE) are shown in Table 1. The addition of minor amounts of TiB2 and TiC particles increased the σ0.2 and σUTS of the composites at room and elevated temperatures. The σ0.2 of the BC-0.05 and BC-0.1 composites were 11.5% and 25.2% higher than that of the Al alloy at 298 K. The TiB2 and TiC particles simultaneously 3

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Fig. 3. The EBSD maps and grain size distribution maps of (a and d) the Al–Cu–Mg–Si alloy, (b and e) the BC-0.05 composite, and (c and f) the BC-0.1 composite. The deformation, substructure and recrystallization regions are marked with red, yellow and blue, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

Fig. 4. (a), (b), and (c) are the tensile engineering stress-strain curves of the Al–Cu–Mg–Si alloy and the dual-phased and bimodal-sized (TiB2–TiC)/Al–Cu–Mg–Si composites at 298 K, 453 K, and 493 K, respectively; (j) the work-hardening rate vs. true strain at 493 K. 4

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Table 1 Details of the tensile tests performed on the Al–Cu–Mg–Si alloy and (TiB2–TiC)/Al–Cu–Mg–Si composites. Samples

298 K

453 K

493 K

σ0.2/MPa

σUTS/MPa

UE/%

σ0.2/MPa

σUTS/MPa

UE/%

σ0.2/MPa

Al alloy

6 278+ −7

10 435+ −9

2.2 19. 0+ −1.0

8 263+ −6

11 357+ −8

0.5 16. 5+ −0.3

BC-0.05

8 310+ −11

7 464+ −10

0.6 22. 8+ −0.3

7 289+ −8

9 383+ −10

0.3 16. 7+ −0.4

BC-0.1

10 348+ −7

8 517+ −9

0.6 5+ −0.4

9 310+ −7

8 413+ −9

0.6 4+ −0.4

17.

17.

σUTS/MPa

UE/%

4 255+ −3

8 318+ −7

0.5 8. 5+ −0.4

5 267+ −4

7 340+ −4

0.7 9. 5+ −0.8

6 279+ −5

4 366+ −3

0.4 10. 6+ −0.5

treatment. It can be inferred that the smaller θ′ and Q′ precipitates in the composite play a more significant role in hindering dislocation motion and increasing the strength of the composite than the larger precipitates. Fig. 8(a) shows the TEM micrograph of the BC-0.1 composite. The TiB2 and TiC particles were located inside the α-Al grains and at the GBs. The dislocation motions were hindered by the particles. The dislocations accumulated and tangled near the particles as the dislocation walls formed. Fig. 8(b) and (c) shows the TEM images and corresponding selected-area electron diffraction (SAED) patterns of the submicron-sized spherical TiC and hexagonal TiB2 particles, respectively, which distributed at the GBs. Fig. 8(d) and (g) show the nanosized cubic TiB2 and spherical TiC particles, respectively. The TiB2 particles exhibited hexagonal and cubic shapes, while the TiC particles had spherical or nearly spherical shapes, which were consistent with those reported in the literature [13,25–29]. Fig. 8(e) shows the HRTEM image of area A in Fig. 8(d), suggesting a clean and well-bonded TiB2/ Al interface. The lattice mismatch between the nano-TiB2 particles and the Al matrix was 3.4%. Fig. 8(f) presents the inverse fast-Fourier-filtered (IFFT) image of area B in Fig. 8(e). The dislocations in the Al matrix were labeled with “T” shaped symbols. A high density of dislocations and severe lattice distortions were observed near the TiB2/Al interface. Fig. 8(h) shows the HRTEM image of area C in Fig. 8(g),

before and after tensile testing at 493 K. The precipitates are observed along the direction parallel to the [001]α-Al axis. As observed in the inset in Fig. 6(d), the Q′ precipitates were coherent with the Al matrix. The mean sizes of the θ′ and Q′ precipitates in the Al alloy were 48 nm and 6.0 nm, respectively, while the mean sizes of the θ′ and Q′ precipitates in the BC-0.1 composite were 34 nm and 3.3 nm, respectively. After tensile testing at 493 K, the diameter of the θ′ precipitates in the Al matrix alloy increased from 48 nm to 94 nm, while that in the composite increased from 34 nm to 70 nm. The sizes of the Q′ precipitates increased from 6.0 nm to 11.4 nm in the Al matrix alloy but increased from 3.3 nm to 7.1 nm in the composite. The composite possessed smaller sizes and higher densities of θ′ and Q′ precipitates than the matrix alloy before and after the elevated-temperature tensile test. As abovementioned, the TiB2 and TiC particles effectively refined the α-Al grains in the as-cast microstructure. The eutectic Al2Cu and AlCuMgSi phases, which were mainly distributed at the grain boundaries, distributed more uniformly in the composites with refined grains, as observed in Fig. 2(d–f). The diffusion distance of the Cu, Mg, and Si atoms in the composites was shorter than that in the Al matrix alloy during the solution treatment [22–24]. The composites obtained a more uniform distribution of Cu, Mg, and Si atoms after the solution treatment. Therefore, a large number of θ′ and Q′ precipitates nucleated simultaneously with as the grain sizes decreased during the aging

Fig. 5. The comparison of the tensile properties of reported Al–Cu metal matrix composites at 493 K, (a) ultimate tensile strength vs. particle content, (b) uniform elongation vs. particle content, and (c) ultimate tensile strength vs. uniform elongation. The volume fraction of the reinforcements was converted to the mass fraction according to Ref. [20] at 473 K and Ref. [21] at 463 K [5,6,8,9,13,18–21].

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Fig. 6. The TEM micrographs of the Al alloy and BC-0.1 composite before tensile testing at 493 K: the θ′ of the (a) Al alloy and (b) BC-0.1 composite, the Q′ of the (c) Al alloy and (d) BC-0.1 composite; (e–h) the corresponding size distribution maps of the θ′ and Q′ precipitates in (a–d), respectively. The inset in Fig. 4(d) is the HRTEM image of zone A.

the dislocation motions and increased the strength of the composites. In terms of the Al alloys, the sliding of GBs was the dominant deformation mechanism at elevated temperatures (0.2–0.5Tm), and the fracture involved the formation, growth, and coalescence of the voids located at the GBs [5,18]. The submicron-sized TiB2 and TiC particles with strong interfacial bonding with the Al matrix were effectively pinned at the grain boundaries, which impeded grain boundary sliding and migration, delayed the formation of microvoids, and helped to increase the elevated-temperature strength and ductility of the composites. In terms of the refined partial recrystallization microstructure of the composites, the fine recrystallized grains contributed to increasing the work hardening capacity, strength, and uniform elongation. Finally, the composites exhibited an excellent elevated temperature strength and uniform elongation.

displaying good TiC/Al interfacial bonding. The lattice mismatch between the nanosized TiC particles and the Al matrix was 5.6%. The lattice distortion near the TiC/Al interface was clearly observed in Fig. 8(i). Due to the thermal expansion mismatch between the TiB2 and TiC particles and the Al matrix, stress occurred near the TiB2/Al and TiC/Al interfaces. With an increase in the number of TiB2 and TiC particles, the dislocation density increased, which promoted the recrystallization of nuclei during hot extrusion [30]. The growth of the equiaxed fine grains was restrained, and the stability of the grains was improved, which was attributed to pinning effects of the submicronsized TiB2 and TiC particles. Thus, the BC-0.05 and BC-0.1 composites had refined recrystallized grains. It is well known that the θ′ and Q′ precipitates hinder dislocation motions and increase the threshold stress for deforming the matrix and composites [11,31]. Although the precipitates grew and coarsened at 493 K, the θ′ and Q′ precipitates in the composite still possessed fine average sizes of 70 nm and 7.1 nm, respectively. The higher density of smaller sized θ′ and Q′ precipitates in the composites helped to increase the strength [5,23]. The nanosized TiB2 and TiC particles also impeded

4. Conclusion Adding minor amounts of dual-phased and submicron/nanosized TiB2–TiC particles simultaneously increased the yield strength (σ0.2),

Fig. 7. TEM micrographs of the Al alloy and BC-0.1 composite after tensile testing at 493 K: the θ′ of the (a) Al alloy and (b) BC-0.1 composite, the Q′ of the (c) Al alloy and (d) BC-0.1 composite; (e–h) the corresponding size distribution maps of the θ′ and Q′ precipitates in (a–d), respectively. 6

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Fig. 8. The TEM micrographs of the BC-0.1 composite after tensile testing at 298 K, (a) microstructure micrograph, (b) submicron-TiC particle, (c) submicron-TiB2 particle, (d) nano-TiB2 particle, (e) HRTEM of zone A, (f) IFFT image of zone B, (g) nano-TiC particle, (h) HRTEM of zone C, and (i) IFFT image of zone D. The insets in Fig. 8(b), (c), (d), and (g) are the corresponding SAED patterns of the particles.

Acknowledgements

ultimate tensile strength (σUTS), and uniform elongation (UE) of the (TiB2–TiC)/Al–Cu–Mg–Si composites at elevated temperatures. The σ0.2, σUTS, and UE of the 0.1 wt% (TiB2–TiC)/Al–Cu–Mg–Si composite were 9.4%, 15.1%, and 24.7% higher than those of the Al alloy at 493 K. The simultaneous increases in the elevated-temperature strength and ductility of the composites were attributed to (i) nanosized TiB2 and TiC particles and the larger density of fine θ′ and Q′ precipitates, hindering dislocation motions, (ii) submicron-sized TiB2 and TiC particles, which were pinned at grain boundaries and impeded grain boundaries sliding, and (iii) the refined partial recrystallization microstructure, which provided an additional contribution to improving the strength and ductility. Compared with the reported Al–Cu matrix composites, the 0.1 wt% (TiB2–TiC)/Al–Cu–Mg–Si composite exhibited the best comprehensive tensile strength and uniform elongation at 493 K. This study provided a new strategy for preparing Al–Cu(-Mg-Si) alloys with superior elevated-temperature strength and ductility by a casting method.

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