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Correlation between microstructure and deformation mechanism in Ti66Nb13Cu8Ni6.8Al6.2 composites at ambient and elevated temperatures
Materials Science & Engineering A 767 (2019) 138448
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Correlation between microstructure and deformation mechanism in Ti66Nb13Cu8Ni6.8Al6.2 composites at ambient and elevated temperatures L.H. Liu a, C. Yang a, *, W.W. Zhang a, Z.Y. Xiao a, L.C. Zhang a, b a b
National Engineering Research Center of Near-net-shape Forming for Metallic Materials, South China University of Technology, Guangzhou, 510640, China School of Engineering, Edith Cowan University, 270 Joondalup Drive, Joondalup, Perth, WA, 6027, Australia
A R T I C L E I N F O
A B S T R A C T
Keywords: Titanium alloys Spark plasma sintering Microstructure Plasticity
In this study, Ti66Nb13Cu8Ni6.8Al6.2 composites with equiaxed and Widmanst€ atten microstructure matrixes were prepared by sintering amorphous and crystalline powder, respectively. The alloy with Widmanst€ atten structure exhibits lower yield and fracture strengths at ambient temperature but higher values at elevated temperatures compared with that with equiaxed structure. Microscopic deformation mechanisms of the alloys with different microstructures were investigated by scanning electron microscope (SEM), transmission electron microscope (TEM) and numerical simulation. The results indicate that the profuse dislocation pile-ups and formation of nanotwins in the β-Ti matrix contribute to the high room-temperature strength and plasticity of the composite with equiaxed β-Ti structure matrix, while stress concentration and cracking in the interior of acicular (Cu, Ni)Ti2 phases in Widmanst€ atten structure matrix deteriorates plastic deformation capacity of the alloy. The higher strength and plasticity for the composite with Widmanst€ atten structure matrix at elevated temperature is ascribed to dislocation pile-ups around grain boundaries of β-Ti and acicular (Cu, Ni)Ti2 phases and the curving of acicular (Cu, Ni)Ti2 phases inside Widmanst€ atten structure matrix. The results obtained provide some guidelines to design high-temperature alloy with excellent properties for structural applications.
1. Introduction Over the past decade, improving the ductility has been a major goal for nanostructured materials [1]. One important strategy was to intro duce some ductile micrometer-sized crystalline phases into nano structured matrix to form a bimodal composite structure [2,3]. Up to now, a lot of Ti-based [4–11], Fe-based [12] and Al-based [13,14] alloys containing micron-scale dendrites in an ultrafine/nano eutectic matrix have been synthesized successfully by rapid solidification. Among them, Ti–Nb/Ta–Cu–Ni/Co–Al alloy system has been received wide concern because of the high strength and large plasticity [4–10].For instance, Okulov et al. reported that the fracture strength and tensile plasticity of bimodal titanium alloys [15–19], such as Ti68.8Nb13.6Co6Cu5.1Al6.5 and Ti68.8Nb13.6Cr5.1Co6Al6.5 alloy [15,16], are always larger than 1.2 GPa and 10%, respectively. However, the size of these alloys fabricated by rapid solidification is so small that hard to meet the requirements for engineering applications. Recently, a material forming method that coupled sintering and crystallizing amorphous powder was carried out to prepare the alloys with larger size [20–27]. The as-sintered Ti–Nb/ Ta–Cu–Ni/Co–Al alloys display an ultrafine and equiaxed composite
microstructure, and possess higher strength and larger plasticity at room temperature compared with the counterparts fabricated by rapid solid ification [24,25]. In the Ti-based bimodal composites fabricated by rapid solidification, the excellent combination of high strength and large plasticity is ascribed to the special microscopic deformation mechanism, such as the rotational boundary [28], interactions of slip bands [28,29], and dislocation multiplication in the micron-scale dendrites [2,29,30]. In contrast, for the ultrafine-grained equiaxed Ti–Nb/Ta–Cu–Ni/Co–Al composites prepared by crystallizing amorphous powder, the origins of higher strength and larger plasticity are still ambiguous. In addition, the most current studies for Ti–Nb/Ta–Cu–Ni/Co–Al composites mainly focus on the mechanical properties at room tem perature [4–15]. The microstructures and properties for these alloys at €tten elevated temperature are received little attention. Widmansta structure, as a characteristic of long nickel-iron crystals, has long been observed in steel [31]. This structure usually makes up of fine inter leaving acicular phases. Generally, the appearance of the Widmanst€ at ten structure in steel dramatically deteriorates the mechanical properties including impact toughness. In the titanium alloys consoli dated by using traditional crystalline powders as precursor, the
* Corresponding author. E-mail address: [email protected] (C. Yang). https://doi.org/10.1016/j.msea.2019.138448 Received 25 March 2019; Received in revised form 29 July 2019; Accepted 22 September 2019 Available online 23 September 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Widmanst€ atten structure is also a common morphology [32,33]. How €tten structures in the Ti–Nb/Ta–Cu–Ni/Co–Al al ever, the Widmansta loys were rarely reported and studied. Meanwhile, the mechanical properties of Ti–Nb/Ta–Cu–Ni/Co–Al alloys at high temperature are not €tten structure. clear, especially for the ones having Widmansta In this study, Ti66Nb13Cu8Ni6.8Al6.2 alloys with equiaxed and Wid €tten microstructure were prepared successfully by sintering mansta amorphous and crystalline powder, respectively. The mechanical property and deformation mechanism for the as-sintered alloys at ambient and elevated temperatures were investigated by SEM, TEM and numerical simulation. The experimental results indicates that the alloy €tten structure matrix displays lower yield strength, with Widmansta fracture strength and plasticity at ambient temperature but higher yield strength and fracture strength at elevated temperature compared with the composite with equiaxed structure matrix. Deformation mechanism analyses show dislocation pile-ups induced by second phases and for mation of nanotwins in the β-Ti matrix contribute to the high strength and large plasticity of alloy with equiaxed structure at room tempera ture. Acicular (Cu, Ni)Ti2 phases in Widmanst€ atten structure matrix deteriorate mechanical property at room temperature due to stress concentration and formation of cracks, but improve the ultimate strength at elevated temperature by curving of acicular phases. The obtained results can give some valuable information on designing alloys with excellent combined mechanical properties at room temperature and elevated temperature.
at elevated temperatures, respectively. The results indicate that the al loys sintered by the with different milling times have the same phase constitutions, i.e., fcc (Cu, Ni)Ti2 phase and bcc β-Ti phase (Fig. 1a, c), which are the same to those in Ti66Nb13Cu8Ni6.8Al6.2 nanostructuredendrites bimodal composite fabricated by rapid solidification [8]. Additionally, it is found that there are not phase transformation after deformation behavior at the elevated temperatures (Fig. 1b, d). The phase constitution of Ti66Nb13Cu8Ni6.8Al6.2 alloys after high tempera ture compressive tests is the same with that of the as-sintered alloys. Fig. 2 shows SEM and TEM images of the Ti66Nb13Cu8Ni6.8Al6.2 bulk alloys fabricated by sintering amorphous and crystalline powder, respectively. Although the alloys have the same phase constitutions (Fig. 1), the phase morphology and distribution of the two alloys are completely different. The sample prepared by sintering and crystallizing amorphous powders displayed with the microstructure of equiaxed fcc (Cu,Ni)Ti2 phases with a grain size of 0.5–1 μm surrounded by equiaxed bcc β-Ti matrix (Fig. 1a), while the alloy fabricated by sintering of crystalline powder displays a Widmanst€ atten structure matrix sur rounded by equiaxed fcc (Cu, Ni)Ti2 phase with a grain size less than 1 μm (Fig. 2b–c). Further bright-field TEM image shows that the Wid manst€ atten structure matrix contains interleaving acicular (Cu, Ni)Ti2 phases surrounded by a bcc β-Ti phases (Fig. 2d). The thickness of the acicular (Cu, Ni)Ti2 regions is about 40 nm, and the interleaving space of the acicular (Cu,Ni)–Ti2 phase is approximately 100 nm (Fig. 2d). The mechanical alloying is a process involving repetitive cold welding and fracture of powders [21,34]. It was found that the as-milled Ti66Nb13Cu8Ni6.8Al6.2 powders always have a layered structure con taining Nb regions and Ti regions [21], which are induced by the cold €tten structure can be formed welding. It was found that the Widmansta by chemical reaction of the layered structure during sintering process [31]. During mechanical alloying process, the thickness of layer struc ture in alloy powders gradually decreases with increasing milling time. Eventually, the alloy elements will distribute evenly in the powders when the glassy phase is formed [35]. The formation of equiaxed structure in the alloys consolidated by glassy powders is mainly ascribed to property of amorphous alloy and characteristics of spark plasma sintering. On the one hand, the (CuNi)Ti2 phase and β-Ti phases are precipitated from amorphous powders. This process involves nucleation and growth [22,36]. The high heating rate of spark plasma sintering contributes to the formation of equiaxed structure [36]. On the other hand, the as-milled glassy powder always possesses highly dense random-packed structure, which would causes the lower diffusion rate during sintering [25,36]. Therefore, the alloy prepared by sintering and
2. Experimental Firstly, crystalline and amorphous Ti66Nb13Cu8Ni6.8Al6.2 (at.%) alloy powders were prepared by mechanical alloying in a high-energy plan etary ball mill (QM-2SP20, apparatus factory of Nanjing University) under a purified argon gas atmosphere. Subsequently, the crystalline (10 h milled) and amorphous powder (50 h milled) were sintered under an argon atmosphere in a Dr. Sintering SPS-825 system with sintering temperature of 900 � C, heating rate of 170 � C/min, holding time of 5 min and pressure of 50 MPa. The diameter of the consolidated sample was 10 mm and the height was 12 mm. Microstructure characterizations were carried out in scanning electron microscopy (SEM; Philips XL-30 FEG, Amsterdam, Netherlands), transmission electron microscopy (TEM; FEI Tecnai G2 F30, Eindhoven, Netherlands) and X-ray diffraction (XRD) (D/MAX-2500/PC; Rigaku Corp., Tokyo, Japan). Compressive tests were performed at room temperature and at elevated temperatures (400 � C, 500 � C and 600 � C) respectively using an MTS universal testing system and a Gleeble 3500 thermal simulation test machine, respec tively. The heating rate in high temperature compression test is 10 K/s and holding time is 30s. The strain rate of the compressive tests under room temperature and at elevated temperatures is 5.0 � 10 4 s 1.The tested samples at elevated temperatures were water quenched imme diately after unloaded in order to retain their high-temperature micro structures. For comparison, a commercial Ti-alloy TG6 with a composition Ti-5.8Al-4.0Sn-4.0Zr-0.7Nb-1.5Ta-0.4Si-0.06C was also tested [33]. Finally, finite element analysis was conducted by using ABAQUS 6.14 software to reveal the microscopic deformation mecha nism of the as-prepared alloys. The Mises stress, which is usually used to determine if a given material will yield or fracture, is calculated by using the equation: σ ¼ ð1=2ððσ1 σ 2 Þ2 þ ðσ2 σ3 Þ2 þ ðσ 3 σ 1 Þ2 ÞÞ1=2 , where σ 1, σ2 and σ3 is principal stress under different direction. The Young’s modulus of β-Ti phases and (Cu, Ni)Ti2 phase is 121 GPa and 42 GPa, respectively. The Poisson’s ratios for both phases were set as 0.3. These values were determined by nanoindentation experiment. The formation of microcrack was not considered in numerical simulation process. 3. Results
Fig. 1. XRD patterns of the Ti66Nb13Cu8Ni6.8Al6.2 sintered by using different powders and after deformation at elevated temperatures: (a) crystalline pow ders; (b) crystalline powder and after deformation at 500 � C; (c) amorphous powders; (b) amorphous powders and after deformation at 500 � C.
Fig. 1 displays the XRD pattern of Ti66Nb13Cu8Ni6.8Al6.2 alloys sin tered by using amorphous and crystalline powder, and after deformation 2
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Table 1 Compressive testing properties for the as-fabricated alloys at ambient and at elevated temperatures. The yield stress σy, ultimate stress σf, fracture strain εf. Herein “Alloy E” and “Alloy W00 denote the alloys with equiaxed and with Widmanst€ atten structures, respectively. Sample
T (oC)
σ y (MPa)
σf (MPa)
εf (%)
Alloy W
20 400 500 600
1380 � 15 1365 � 12 1260 � 16 610 � 10
1885 � 11 1500 � 16 1153 � 13 300 � 12
13.3 � 0.2 24.1 � 0.1 23.2 � 0.3 >50
Alloy E
20 400 500 600
1450 � 12 1360 � 8 853 � 15 400 � 14
2350 � 11 1405 � 13 869 � 13 239 � 12
28.5 � 0.2 35.2 � 0.1 >50 >50
TG6
400 500 600
950 � 9 697 � 13 650 � 12
905 � 24 700 � 12 550 � 22
20.4 � 0.1 27.5 � 0.3 >50
equiaxed structure are also larger than other advanced titanium alloys fabricated by melt solidification, including bulk metallic glass [37–39] and nanostructure-dendrite bimodal composites [4–8,16–18]. In €tten structure matrix exhibits higher contrast, the alloy with Widmansta yield strength and ultimate strength at elevated temperatures than those of the alloy with equiaxed bcc β-Ti matrix (Fig. 3c–d and Table 1). The yield strength, fracture strength and fracture strain of the alloy with €tten microstructure matrix at 500 � C is 1260 MPa, 1153 MPa, Widmansta and 23.2%, respectively (Fig. 3c), which are superior to those of the commercial Ti-alloy TG6. Besides, it is found that the yield strength of the alloy W decreases drastically as the testing temperature is increased to 600 � C, while this softening temperature point decreases to about 500 � C for the alloy E (Table 1). The results indicate the alloy W with €tten structure matrix have a better high temperature Widmansta resistance. In order to clarify origins of different mechanical properties for the as-fabricated alloys at room and elevated temperatures, TEM in vestigations were conducted on the specimens deformed (8%-strained) under different temperatures. As shown in Fig. 4a, nano-sized cracks arise at the grain boundaries between the (Cu, Ni)Ti2 phases and be tween the (Cu, Ni)Ti2 and β-Ti matrix in the alloy E. Meanwhile, abundant dislocations were clearly observed in the β-Ti matrix and hindered by the (Cu, Ni)Ti2 phases, but seldom were found in (Cu,Ni)Ti2 phases. The results indicate that the β-Ti phase regions have the excel lent plastic deformation and work-hardening capacity thereby acts as the ductile phases with respect to the (Cu, Ni)Ti2 phases. Besides, TEM observations show that there are lots of deformation-induced nanotwins (or stacking faults) in the equiaxed β-Ti matrix in alloy E (Fig. 4b). The width of the nanostructure determined by HRTEM is about eight inter atomic distances. The corresponding SAED pattern of the nanostructure exhibits elongated diffraction spots along the crystal planes (110) and (1ð Þ10) (Fig. 4b inset), indicating that there are severe lattice dis tortions inside this region. In contrast, the 8%-strained at 500 � C and then quenched sample exhibits completely different morphologies from the corresponding counterparts deformed at room temperature (Fig. 4c). Firstly, the equiaxed fcc (Cu, Ni)Ti2 phases in alloy E undergoes distinct plastic deformation at high temperature. This can be testified by their elliptic or irregular shapes as indicated by the red dashed lines (Fig. 4c). Secondly, no micro-cracks and damaged interface between the (Cu, Ni) Ti2 and β-Ti phases were observed in the high-temperature deformed sample (Fig. 4d), different from that formation of nano-sized cracks in the room-temperature deformed sample (Fig. 4a). The results indicates that the plastic deformation ability of (Cu, Ni)Ti2 phases increases drastically with increasing of testing temperature. The microstructures for the deformed Ti66Nb13Cu8Ni6.8Al6.2 com posite with Widmanst€ atten structure matrix are displayed in Fig. 5. Similar to the sample with equiaxed bcc β-Ti matrix, deformationinduced dislocations and cracking were formed in the β-Ti phases and
Fig. 2. SEM microstructures of the Ti66Nb13Cu8Ni6.8Al6.2 bulk alloys fabricated by sintering from (a) amorphous and (b) crystalline powder. (c) TEM image and (d) corresponding magnified Widmanst€ atten structure matrix of the alloy in (b).
crystallizing amorphous powders always displays the equiaxed struc ture, while the alloy fabricated by sintering of crystalline powder pos €tten structure. sesses a Widmansta Fig. 3 presents the compressive engineering stress-strain curves of the sintered Ti66Nb13Cu8Ni6.8Al6.2 bulk alloys with equiaxed bcc β-Ti €tten matrix (alloy W) at room matrix (alloy E) and with Widmansta temperature and at elevated temperature. For comparison, the stressstrain curves of commercial high-temperature titanium alloy TG6 is also displayed in Fig. 3. The corresponding mechanical properties are listed in Table 1. With increasing of test temperature, the yield strength and fracture strength decrease for both alloy E and alloy W (Table 1). Meanwhile, it is found that the alloy with equiaxed structure matrix has higher yield strength and ultimate strength and larger plasticity at room €tten structure ma temperature than those of the alloy with Widmansta trix (Fig. 3a–b and Table 1). The strength and plasticity in the alloy with
Fig. 3. Compressive engineering stress-strain curves at room temperature (a and b) and elevated temperature (c and d) of the fabricated Ti66Nb13Cu8 Ni6.8Al6.2 bulk alloys with different structures. Herein “E” and “W” denote the alloys with equiaxed and with Widmanst€ atten structure matrix, respectively. (e) The commercial high-temperature alloy TG6. 3
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plasticity in room-temperature. When the samples is deformed at 500 � C, interestingly, it can be found that the acicular (Cu, Ni)Ti2 phases are curved into the “C” or “S” shapes (Fig. 5c–d). Meanwhile, profuse dislocations are hindered and piled up in the β-Ti matrix where is close to the (Cu, Ni)Ti2 phases, showing that dislocations are hard to cross the boundary of β-Ti phase and (Cu, Ni)Ti2 phases. The corresponding SAED pattern confirms that “C” or “S” shaped phases possess fcc structure and the high temperature does not cause phase transformation of the (Cu,Ni) Ti2 phase (Fig. 5d), which is consistent with the XRD results (Fig. 1). Recently, finite element analysis (FMA) as a supplementary method was applied to reveal the essential relationship of microstructure and property of metallic materials. For instance, Jeon et al. employed the FMA method to simulate deformation behaviors of Ti-based amorphous matrix composites [40]. The FMA results of deformation bands or shear bands are in good agreement with the experimental findings [40]. In order to investigate further the effects of microstructure on the ductility of the present Ti66Nb13Cu8Ni6.8Al6.2 composites, the FMA was per formed to reveal the microscopic deformation mechanism. According to the above SEM and TEM observations, it can be concluded that the €tten structure matrix and equi biggest difference between Widmansta axed structure matrix is that there are acicular (Cu,Ni)Ti2 phases in the €tten structure but not in equiaxed structure. Therefore, the Widmansta numerical simulation models for the alloys with different microstructure are simplified as Fig. 6a and Fig. 6c, respectively. The white regions denote (Cu,Ni)Ti2 phases and blue regions donate the β-Ti phase. The Mises stress fields in the composites under different strains are displayed in Fig. 6b–c and Fig. 6e-f, respectively. When the strain is 0.6%, the maximum stress in the alloy with equiaxed β-Ti phase matrix is mainly distributed at junction of the (Cu,Ni)Ti2 phases and (Cu,Ni)Ti2 phases, or at interface between the (Cu,Ni)Ti2 phase and the matrix (Fig. 6b). The stress concentration increases with increase of strain (Fig. 6c). In the samples having acicular (Cu,Ni)Ti2 phases, the maximum stress also distributes at the boundary of the (Cu,Ni)Ti2 phases when the strain is small (Fig. 6e). However, the stress begins to concentrate in the middle of acicular (Cu,Ni)Ti2 when the sample is yielded (Fig. 6f). The nu merical simulation results indicate that acicular (Cu,Ni)Ti2 phases in the alloys play a critical key role on the stress distribution during deformation. The evolutions of Mises stress values extracted from the different positions in the samples are displayed in Fig. 7. During the elastic deformation stage, the stress values in the β-Ti phases are higher than those both in equiaxed and in acicular (Cu,Ni)Ti2 phases (Fig. 7a–b), which is mainly ascribed to the fact that elastic modulus of β-Ti phases is larger than that of (Cu,Ni)Ti2 phases [36]. With increase of loading, in alloy E with equiaxed structure the Mises stress at the interface between the (Cu,Ni)Ti2 phase and β-phase matrix increase sharply (Fig. 7a). When the strain is 8%, the maximum stress at the junction reaches about 2023 MPa, which is about 1.6 times of loading (Fig. 7a). The result in dicates that cracking prefers to form at interface between the (Cu,Ni)Ti2 phase and β-phase matrix in the alloy E with equiaxed structure, which agrees well with the TEM observations (Fig. 4a). In the sample W with Widmanst€ atten structure, Mises stress at interface between the (Cu,Ni) Ti2 phase and β-phase matrix is also larger than that in matrix or second phases. However, the Mises stress values in the middle of acicular (Cu, Ni)Ti2 phases increases rapidly (Fig. 7b). When the strain is 8%, the maximum values in the acicular (Cu,Ni)Ti2 phases can reach 3701 MPa, which is far larger than that at the interface of (Cu,Ni)Ti2 phases and β-phase matrix. The result indicates the cracks in the samples with Widmanstδtten structure will be firstly formed in the middle position of acicular (Cu,Ni)Ti2 phases.
Fig. 4. (a) TEM bright-field micrograph of the 8%-strained sample E with equiaxed bcc β-Ti matrix at room temperature, (b) Deformation induced nanotwins in β-Ti matrix. The insets are corresponding SAED pattern and highresolution TEM (HRTEM) micrograph of the square region respectively. (c) TEM bright-field micrograph of the 8%-strained sample E with equiaxed bcc β-Ti matrix at 500 � C, (d) High-resolution TEM image of the phase boundary of the β-Ti matrix and fcc (Cu, Ni)Ti2 phase in (c).
at the boundaries between the equiaxed second phase and matrix (Fig. 5a), respectively. However, a large number of micro-cracks was formed in the interior of acicular (Cu,Ni)Ti2 phases (Fig. 5b). The results indicate that acicular (Cu,Ni)Ti2 phases may be the origin of the low
4. Discussion
Fig. 5. (a) TEM bright-field micrograph of the 8%-strained sample with Wid manst€ atten structure matrix at room temperature, (b) micro-cracks in middle of acicular (Cu, Ni)Ti2 phase. (c–d) TEM morphologies of the acicular (Cu, Ni)Ti2 phases in the 8%-strained sample at 500 � C. The inset in (d) is the corre sponding SAED pattern of fcc (Cu, Ni)Ti2 phase.
4.1. Deformation at room temperature In the previous works, lots of equiaxed structural Ti–Nb/Ta–Cu–Ni/ 4
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Fig. 6. Finite element analysis models used for calculation of Ti-based composites with equiaxed structure (a) and Widmanst€ atten structure (d); Stress field of the equiaxed structural composites applied 0.6% strain (b) and 8% strain (c); Stress field of the Widmanst€ atten structural composites after being applied to 0.6% strain (e) and 8% strain (f).
Co–Al composites have been prepared successfully by crystallizing amorphous powder [20–27], but the origins of high strength and large plasticity are ambiguous. Although the microstructures including pile-up of dislocations and cracks at the interface between the TiFe phase and the matrix have been observed in deformed samples by Li et al. [41,42], up to now the detailed microscopic deformation mecha nism is not elaborated clearly, especially in the samples having acicular phases. In this work, the evolutions of Mises stress during deformation are revealed by FMA method (Fig. 7). By coupling with TEM results, the deformation mechanism of the present Ti66Nb13Cu8Ni6.8Al6.2 compos ites is proposed. The corresponding schematic diagram is displayed in Fig. 8. When loading is small, in the sample with equiaxed bcc β-Ti matrix, elastic deformation occurs simultaneously in both the β-Ti matrix and the (Cu,Ni)Ti2 phases, but the β-Ti phases present considerably higher stress than the (Cu, Ni)Ti2 phases due to the lower modulus of (Cu, Ni) Ti2 phases (Fig. 8a, corresponding to I stage in Fig. 7a). With increasing of stress beyond the yield point of the alloy, a large number of disloca tions are generated in the ductile β-Ti matrix and are pinned by the surrounded equiaxed (Cu, Ni)Ti2 phases (Fig. 8b, corresponding to II stage in Fig. 7a). The profuse dislocation pile-ups would largely contribute to the increase in strength due to working hardening effect [18]. With increasing further the loading, stress concentration increases drastically at the boundaries of equiaxed (Cu, Ni)Ti2 phases (Fig. 7a). When the stress level is increased to a critical value, cracks as shown in Fig. 4a would be generated at boundaries of the equiaxed (Cu, Ni)Ti2 phase regions (Fig. 8c, corresponding to the III stage in Fig. 7a). Since the (Cu, Ni)Ti2 phase in the Ti66Nb13Cu8Ni6.8Al6.2 alloys represents considerably lower flow stress and surrounded by β-Ti phases, the deformation of the (Cu, Ni)Ti2 phase is constrained by the β-Ti phase, consequently, the formed cracks at the boundary are separated and restricted by the adjacent ductile β-Ti matrix (Fig. 4a). Therefore, they are unable to shear-off the whole sample [29]. With increase of stress, the sustaining retention of the increasing number in dislocation pile-ups and strain induces formation of nanotwins in the ductile β-Ti matrix, as
shown in Fig. 4b. It was reported that twinning after an initial dislocation-meditated hardening process could allow the bcc β-Ti phase to undergo superplastic-like deformation [42], which would promote the increases of the plastic strain of samples (Fig. 8d, corresponding to the IV stage in Fig. 7a). When the load is increased to a high enough value that β-Ti matrix cannot suppress propagation of micro-cracks, the cracks would pass through the whole matrix, thus leading to the macroscopic fracture of the as-fabricated bulk alloy. Therefore, the equiaxed Ti66Nb13Cu8Ni6.8Al6.2 composites prepared by crystallizing amorphous powders have high fracture strength and large plasticity, and the excellent mechanical properties mainly originate from the special microstructure and corresponding deformation mechanism. Similar microscopic deformation process arises in the alloy with €tten structure matrix. The only difference is deformation Widmansta behavior of acicular (Cu, Ni)Ti2 phases. It was recognized that hardness of β-Ti phase was three times that of (Cu, Ni)Ti2 phase [36]. This means that the strength of (Cu, Ni)Ti2 phase is lower than that of β-Ti phase, and thus (Cu, Ni)Ti2 phase is easier to break than β-Ti phase. Therefore, €tten structure can the deformation process for the alloy with Widmansta be divided into following several stages. When the loading is small, the stress concentration distributes in the interface between the (Cu,Ni)Ti2 phase and β-phase matrix (Fig. 8a, corresponding to I stage in Fig. 7b). With increasing of loading, stress concentration increases drastically at the interior of acicular (Cu, Ni)Ti2 phases. Subsequently, the cracks would be formed in the acicular (Cu, Ni)Ti2 phases (Fig. 8b, corre sponding to II stage in Fig. 7b). As the stress increases further, larger sized cracks will be formed at the boundary of micron-sized (Cu, Ni)Ti2 phases like that in alloy with equiaxed structure (Fig. 8c, corresponding to III stage in Fig. 7b). Since the density of acicular (Cu, Ni)Ti2 phases in €tten structure matrix is high, the formed cracks in the acic Widmansta ular phase would deteriorate plastic deformation ability of Wid manst€ atten structure matrix. Therefore, once the crack is formed in boundary of micron-sized (Cu, Ni)Ti2 phases, they are hard to be sup pressed by the adjacent Widmanst€ atten structure matrix having micro cracks. As a result, the samples will break quickly with low plasticity and 5
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strength (Fig. 8d, corresponding to IV stage in Fig. 7b). Therefore, the €tten structure matrix ex Ti66Nb13Cu8Ni6.8Al6.2 alloy with Widmansta hibits lower strength and less plasticity than that with equiaxed bcc β-Ti matrix at room temperature. 4.2. Deformation at elevated temperature As shown in Fig. 3 and Table 1, the yield strength of the two samples with equiaxed bcc β-Ti matrix and with Widmanst€ atten structure matrix decreases with increasing the testing temperature. The yield point is always a critical stress to start emission of dislocation or activation of sliding. It is widely recognized that the critical stress decreases with increasing the temperature [43,44]. Besides, the yield strength in the composite is also determined by second phases [45]. TEM observation shows that the equiaxed (Cu, Ni)Ti2 second phases are compressed into elliptic or irregular shapes under deformation, indicating that strength of the (Cu, Ni)Ti2 second phases decreases sharply at the high temper ature. Therefore, the yield strength for each sample decreases with increasing the testing temperature. With regard to the plasticity, it is associated with the deformation mechanism. It has been reported that the glide of grains or phase boundaries becomes easier with increasing temperature [43,44]. TEM observations also demonstrate that the deformation abilities of the (Cu, Ni)Ti2 and β-Ti phases increase with increasing the temperature (Fig. 4c–d and Fig. 5c–d). The increase of plastic deformation capacity will lead to release of the stress concen tration resulting in crack nucleation in grain boundaries. Therefore, the plasticity increased with the increasing test temperature for each specimen. Also, it can be noticed that the sample with Widmanst€ atten structure matrix exhibits higher yield strength and fracture strength than the alloy ^-Ti matrix at elevated temperatures, indicating its with equiaxed bcc a higher deformation resistance ability to elevated temperatures. This is ^-Ti phase content in alloys with Widmansta €tten mainly ascribed to the a structure and deformation mechanism of the acicular (Cu,Ni)Ti2 phases. Under the high temperature, it was reported that the minimum sliding resistance in titanium alloy is the β phase/β phase boundary [46,47], and only the dislocations which slipped along β-Ti slip system can cross the β/α interfaces easily [33]. As shown in Fig. 2d, the β-Ti phases in €tten structure are interleaved by the acicular (Cu, Ni)Ti2 Widmansta phases. Besides, the volume fraction of β-Ti in the alloys with Wid manst€ atten structure matrix is far lower than that in the samples with equiaxed bcc β-Ti matrix. Therefore, the content of β/β boundary in the €tten structure is less than that in alloy with equi alloy with Widmansta axed structure. So the alloy with Widmanst€ atten structure matrix pos sesses higher yield strength at elevated temperatures (Table 1). Also, it can be seen that the equiaxed (Cu, Ni)Ti2 phases are elongated at high temperature (Fig. 4c), while the curving of the acicular (Cu, Ni)Ti2 phases is observed in Widmanst€ atten structure matrix (Fig. 5c–d). This indicates that the acicular (Cu, Ni)Ti2 phases can hinder the €tten higher-temperature flow behavior of the sample with Widmansta matrix thereby result in a higher ultimate strength. Shortly, the presence €tten structure matrix of the acicular (Cu, Ni)Ti2 phases inside Widmansta is the decisive reason for the higher yield and ultimate strength at elevated temperature.
Fig. 7. Mises stress extracted from the numerical simulation results of alloys with different structure: (a) equiaxed structure; (b) Widmanst€ atten structure. The “1”, “2”, “3”, and “4” denotes different positions marked in Fig. 5a and d.
5. Conclusion Deformation mechanism of Ti66Nb13Cu8Ni6.8Al6.2 composites with equiaxed β-Ti and with Widmanst€ atten structure matrix was investi gated at ambient and elevated temperatures. Results show that the equiaxed β-Ti structure matrix has higher capability of plastic defor mation with formed nanotwins, while cracking arises in the interior of acicular (Cu, Ni)Ti2 phases in Widmanst€ atten structure matrix. This leads to the higher strength and plasticity of the sample with equiaxed structure matrix at ambient temperature. In contrast, abundant dislo cations generated in β-Ti phases are hindered by curved acicular (Cu, Ni)
Fig. 8. Schematic diagram of deformation behavior in Ti66Nb13Cu8Ni6.8Al6.2 bulk alloys with two different matrixes at room temperature. The left and right ones in each figure denote the alloy with Widmanst€ atten structure matrix and with equiaxed bcc β-Ti matrix respectively. The “I”, “II”, “III” and “IV” stage correspond to the different regions in Fig. 6. 6
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€tten structure matrix at elevated tempera Ti2 phases inside Widmansta €tten tures. This causes the higher strength of the sample with Widmansta structure matrix. The results obtained would provide some insights into designing high-performance metallic alloys used as high-temperature structural materials.
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Data availability statement The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgments This work was supported by China Postdoctoral Science Foundation (No. 2019TQ0099), the National Natural Science Foundation of China (Nos. 51574128 and 51627805)and the Guangdong Natural Science Foundation for Research Team (No. 2015A030312003). References [1] E. Ma, Eight routes to improve the tensile ductility of bulk nanostructured metals and alloys, J. Occup. Med. 58 (2006) 49–53. https://doi.org/10.1007/s11837-00 6-0215-5. [2] S.W. Lee, J.T. Kim, S.H. Hong, H.J. Park, J. Park, N.S. Lee, Y. Seo, J.Y. Suh, J. Eckert, D.H. Kim, J.M. Park, K.B. Kim, Micro-to-nano-scale deformation mechanisms of a bimodal ultrafine eutectic composite, Sci. Rep. 4 (2014) 6500. https://doi.org/10.1038/srep06500. [3] G. He, J. Eckert, W. L€ oser, L. Schultz, Novel Ti-base nanostructure-dendrite composite with enhanced plasticity, Nat. Mater. 2 (2003) 33–37. https://doi. org/10.1038/nmat792. [4] G. He, J. Eckert, W. L€ oser, M. Hagiwara, Composition dependence of the microstructure and the mechanical properties of nano/ultrafine-structured Ti-CuNi-Sn-Nb alloys, Acta Mater. 52 (2004) 3035–3046. https://doi.org/10.1016/j. actamat.2004.03.006. [5] D.V. Louzguine-Luzgin, L.V. Louzguina-Luzgina, H. Kato, A. Inoue, Investigation of Ti-Fe-Co bulk alloys with high strength and enhanced ductility, Acta Mater. 53 (2005) 2009–2017. https://doi.org/10.1016/j.actamat.2005.01.012. [6] L. Zhang, H. Lu, C. Mickel, J. Eckert, Ductile ultrafine-grained Ti-based alloys with high yield strength, Appl. Phys. Lett. 91 (2007), 051906. https://doi.org/10.1063/ 1.2766861. [7] L.C. Zhang, J. Das, H.B. Lu, C. Duhamel, M. Calin, J. Eckert, High strength Ti-Fe-Sn ultrafine composites with large plasticity, Scr. Mater. 57 (2007) 101–104. https://d oi.org/10.1016/j.scriptamat.2007.03.031. [8] U. Kühn, N. Mattern, A. Gebert, M. Kusy, M. Bostr€ om, U. Siegel, L. Schultz, Nanostructured Zr- and Ti-based composite materials with high strength and enhanced plasticity, J. Appl. Phys. 98 (2005) 54307. https://doi.org/10.1063/ 1.2012507. [9] L.C. Zhang, L.Y. Chen, A review on biomedical titanium alloys: recent progress and prospect, Adv. Eng. Mater. 21 (2019) 1801215. https://doi.org/10.1002/adem. 201801215. [10] I.V. Okulov, U. Kühn, T. Marr, J. Freudenberger, I.V. Soldatov, L. Schultz, C. G. Oertel, W. Skrotzki, J. Eckert, Microstructure and mechanical properties of new composite structured Ti-V-Al-Cu-Ni alloys for spring applications, Mater. Sci. Eng. A 603 (2014) 76–83. https://doi.org/10.1016/j.msea.2014.02.070. [11] C.D. Rabadia, Y.J. Liu, L. Wang, H. Sun, L.C. Zhang, Laves phase precipitation in TiZr-Fe-Cr alloys with high strength and large plasticity, Mater. Des. 154 (2018) 228–238. https://doi.org/10.1016/j.matdes.2018.05.035. [12] J.T. Kim, S.H. Hong, H.J. Park, Y.S. Kim, G.H. Park, J.Y. Park, N. Lee, Y. Seo, J. M. Park, K.B. Kim, Improving the plasticity and strength of Fe-Nb-B ultrafine eutectic composite, Mater. Des. 76 (2015) 190–195. https://doi.org/10.1016/j. matdes.2015.03.053. [13] C.S. Tiwary, D. Roy Mahapatra, K. Chattopadhyay, Effect of length scale on mechanical properties of Al-Cu eutectic alloy, Appl. Phys. Lett. 101 (2012) 171901. https://doi.org/10.1063/1.4761944. [14] J.M. Park, N. Mattern, U. Kühn, J. Eckert, High-strength bulk Al-based bimodal ultrafine eutectic composite with enhanced plasticity, J. Mater. Res. 24 (2009) 2605–2609. https://doi.org/10.1557/jmr.2009.0297. [15] I.V. Okulov, H. Wendrock, A.S. Volegov, H. Attar, U. Kühn, W. Skrotzki, J. Eckert, High strength beta titanium alloys: new design approach, Mater. Sci. Eng. A 628 (2015) 297–302. https://doi.org/10.1016/j.msea.2015.01.073. [16] I.V. Okulov, M. B€ onisch, U. Kühn, W. Skrotzki, J. Eckert, Significant tensile ductility and toughness in an ultrafine-structured Ti68.8Nb13.6Co6Cu5.1Al6.5 bimodal alloy, Mater. Sci. Eng. A 615 (2014) 457–463. https://doi.org/10.1016/j. msea.2014.07.108. [17] I.V. Okulov, U. Kühn, J. Romberg, I.V. Soldatov, J. Freudenberger, L. Schultz, A. Eschke, C.-G. Oertel, W. Skrotzki, J. Eckert, Mechanical behavior and tensile/ compressive strength asymmetry of ultrafine structured Ti-Nb-Ni-Co-Al alloys with bi-modal grain size distribution, Mater. Des. 62 (2014) 14–20. https://doi.org/10 .1016/j.matdes.2014.05.007.
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Correlation between microstructure and deformation mechanism in Ti66Nb13Cu8Ni6.8Al6.2 composites at ambient and elevated temperatures L.H. Liu a, C. Yang a, *, W.W. Zhang a, Z.Y. Xiao a, L.C. Zhang a, b a b
National Engineering Research Center of Near-net-shape Forming for Metallic Materials, South China University of Technology, Guangzhou, 510640, China School of Engineering, Edith Cowan University, 270 Joondalup Drive, Joondalup, Perth, WA, 6027, Australia
A R T I C L E I N F O
A B S T R A C T
Keywords: Titanium alloys Spark plasma sintering Microstructure Plasticity
In this study, Ti66Nb13Cu8Ni6.8Al6.2 composites with equiaxed and Widmanst€ atten microstructure matrixes were prepared by sintering amorphous and crystalline powder, respectively. The alloy with Widmanst€ atten structure exhibits lower yield and fracture strengths at ambient temperature but higher values at elevated temperatures compared with that with equiaxed structure. Microscopic deformation mechanisms of the alloys with different microstructures were investigated by scanning electron microscope (SEM), transmission electron microscope (TEM) and numerical simulation. The results indicate that the profuse dislocation pile-ups and formation of nanotwins in the β-Ti matrix contribute to the high room-temperature strength and plasticity of the composite with equiaxed β-Ti structure matrix, while stress concentration and cracking in the interior of acicular (Cu, Ni)Ti2 phases in Widmanst€ atten structure matrix deteriorates plastic deformation capacity of the alloy. The higher strength and plasticity for the composite with Widmanst€ atten structure matrix at elevated temperature is ascribed to dislocation pile-ups around grain boundaries of β-Ti and acicular (Cu, Ni)Ti2 phases and the curving of acicular (Cu, Ni)Ti2 phases inside Widmanst€ atten structure matrix. The results obtained provide some guidelines to design high-temperature alloy with excellent properties for structural applications.
1. Introduction Over the past decade, improving the ductility has been a major goal for nanostructured materials [1]. One important strategy was to intro duce some ductile micrometer-sized crystalline phases into nano structured matrix to form a bimodal composite structure [2,3]. Up to now, a lot of Ti-based [4–11], Fe-based [12] and Al-based [13,14] alloys containing micron-scale dendrites in an ultrafine/nano eutectic matrix have been synthesized successfully by rapid solidification. Among them, Ti–Nb/Ta–Cu–Ni/Co–Al alloy system has been received wide concern because of the high strength and large plasticity [4–10].For instance, Okulov et al. reported that the fracture strength and tensile plasticity of bimodal titanium alloys [15–19], such as Ti68.8Nb13.6Co6Cu5.1Al6.5 and Ti68.8Nb13.6Cr5.1Co6Al6.5 alloy [15,16], are always larger than 1.2 GPa and 10%, respectively. However, the size of these alloys fabricated by rapid solidification is so small that hard to meet the requirements for engineering applications. Recently, a material forming method that coupled sintering and crystallizing amorphous powder was carried out to prepare the alloys with larger size [20–27]. The as-sintered Ti–Nb/ Ta–Cu–Ni/Co–Al alloys display an ultrafine and equiaxed composite
microstructure, and possess higher strength and larger plasticity at room temperature compared with the counterparts fabricated by rapid solid ification [24,25]. In the Ti-based bimodal composites fabricated by rapid solidification, the excellent combination of high strength and large plasticity is ascribed to the special microscopic deformation mechanism, such as the rotational boundary [28], interactions of slip bands [28,29], and dislocation multiplication in the micron-scale dendrites [2,29,30]. In contrast, for the ultrafine-grained equiaxed Ti–Nb/Ta–Cu–Ni/Co–Al composites prepared by crystallizing amorphous powder, the origins of higher strength and larger plasticity are still ambiguous. In addition, the most current studies for Ti–Nb/Ta–Cu–Ni/Co–Al composites mainly focus on the mechanical properties at room tem perature [4–15]. The microstructures and properties for these alloys at €tten elevated temperature are received little attention. Widmansta structure, as a characteristic of long nickel-iron crystals, has long been observed in steel [31]. This structure usually makes up of fine inter leaving acicular phases. Generally, the appearance of the Widmanst€ at ten structure in steel dramatically deteriorates the mechanical properties including impact toughness. In the titanium alloys consoli dated by using traditional crystalline powders as precursor, the
* Corresponding author. E-mail address: [email protected] (C. Yang). https://doi.org/10.1016/j.msea.2019.138448 Received 25 March 2019; Received in revised form 29 July 2019; Accepted 22 September 2019 Available online 23 September 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Materials Science & Engineering A 767 (2019) 138448
Widmanst€ atten structure is also a common morphology [32,33]. How €tten structures in the Ti–Nb/Ta–Cu–Ni/Co–Al al ever, the Widmansta loys were rarely reported and studied. Meanwhile, the mechanical properties of Ti–Nb/Ta–Cu–Ni/Co–Al alloys at high temperature are not €tten structure. clear, especially for the ones having Widmansta In this study, Ti66Nb13Cu8Ni6.8Al6.2 alloys with equiaxed and Wid €tten microstructure were prepared successfully by sintering mansta amorphous and crystalline powder, respectively. The mechanical property and deformation mechanism for the as-sintered alloys at ambient and elevated temperatures were investigated by SEM, TEM and numerical simulation. The experimental results indicates that the alloy €tten structure matrix displays lower yield strength, with Widmansta fracture strength and plasticity at ambient temperature but higher yield strength and fracture strength at elevated temperature compared with the composite with equiaxed structure matrix. Deformation mechanism analyses show dislocation pile-ups induced by second phases and for mation of nanotwins in the β-Ti matrix contribute to the high strength and large plasticity of alloy with equiaxed structure at room tempera ture. Acicular (Cu, Ni)Ti2 phases in Widmanst€ atten structure matrix deteriorate mechanical property at room temperature due to stress concentration and formation of cracks, but improve the ultimate strength at elevated temperature by curving of acicular phases. The obtained results can give some valuable information on designing alloys with excellent combined mechanical properties at room temperature and elevated temperature.
at elevated temperatures, respectively. The results indicate that the al loys sintered by the with different milling times have the same phase constitutions, i.e., fcc (Cu, Ni)Ti2 phase and bcc β-Ti phase (Fig. 1a, c), which are the same to those in Ti66Nb13Cu8Ni6.8Al6.2 nanostructuredendrites bimodal composite fabricated by rapid solidification [8]. Additionally, it is found that there are not phase transformation after deformation behavior at the elevated temperatures (Fig. 1b, d). The phase constitution of Ti66Nb13Cu8Ni6.8Al6.2 alloys after high tempera ture compressive tests is the same with that of the as-sintered alloys. Fig. 2 shows SEM and TEM images of the Ti66Nb13Cu8Ni6.8Al6.2 bulk alloys fabricated by sintering amorphous and crystalline powder, respectively. Although the alloys have the same phase constitutions (Fig. 1), the phase morphology and distribution of the two alloys are completely different. The sample prepared by sintering and crystallizing amorphous powders displayed with the microstructure of equiaxed fcc (Cu,Ni)Ti2 phases with a grain size of 0.5–1 μm surrounded by equiaxed bcc β-Ti matrix (Fig. 1a), while the alloy fabricated by sintering of crystalline powder displays a Widmanst€ atten structure matrix sur rounded by equiaxed fcc (Cu, Ni)Ti2 phase with a grain size less than 1 μm (Fig. 2b–c). Further bright-field TEM image shows that the Wid manst€ atten structure matrix contains interleaving acicular (Cu, Ni)Ti2 phases surrounded by a bcc β-Ti phases (Fig. 2d). The thickness of the acicular (Cu, Ni)Ti2 regions is about 40 nm, and the interleaving space of the acicular (Cu,Ni)–Ti2 phase is approximately 100 nm (Fig. 2d). The mechanical alloying is a process involving repetitive cold welding and fracture of powders [21,34]. It was found that the as-milled Ti66Nb13Cu8Ni6.8Al6.2 powders always have a layered structure con taining Nb regions and Ti regions [21], which are induced by the cold €tten structure can be formed welding. It was found that the Widmansta by chemical reaction of the layered structure during sintering process [31]. During mechanical alloying process, the thickness of layer struc ture in alloy powders gradually decreases with increasing milling time. Eventually, the alloy elements will distribute evenly in the powders when the glassy phase is formed [35]. The formation of equiaxed structure in the alloys consolidated by glassy powders is mainly ascribed to property of amorphous alloy and characteristics of spark plasma sintering. On the one hand, the (CuNi)Ti2 phase and β-Ti phases are precipitated from amorphous powders. This process involves nucleation and growth [22,36]. The high heating rate of spark plasma sintering contributes to the formation of equiaxed structure [36]. On the other hand, the as-milled glassy powder always possesses highly dense random-packed structure, which would causes the lower diffusion rate during sintering [25,36]. Therefore, the alloy prepared by sintering and
2. Experimental Firstly, crystalline and amorphous Ti66Nb13Cu8Ni6.8Al6.2 (at.%) alloy powders were prepared by mechanical alloying in a high-energy plan etary ball mill (QM-2SP20, apparatus factory of Nanjing University) under a purified argon gas atmosphere. Subsequently, the crystalline (10 h milled) and amorphous powder (50 h milled) were sintered under an argon atmosphere in a Dr. Sintering SPS-825 system with sintering temperature of 900 � C, heating rate of 170 � C/min, holding time of 5 min and pressure of 50 MPa. The diameter of the consolidated sample was 10 mm and the height was 12 mm. Microstructure characterizations were carried out in scanning electron microscopy (SEM; Philips XL-30 FEG, Amsterdam, Netherlands), transmission electron microscopy (TEM; FEI Tecnai G2 F30, Eindhoven, Netherlands) and X-ray diffraction (XRD) (D/MAX-2500/PC; Rigaku Corp., Tokyo, Japan). Compressive tests were performed at room temperature and at elevated temperatures (400 � C, 500 � C and 600 � C) respectively using an MTS universal testing system and a Gleeble 3500 thermal simulation test machine, respec tively. The heating rate in high temperature compression test is 10 K/s and holding time is 30s. The strain rate of the compressive tests under room temperature and at elevated temperatures is 5.0 � 10 4 s 1.The tested samples at elevated temperatures were water quenched imme diately after unloaded in order to retain their high-temperature micro structures. For comparison, a commercial Ti-alloy TG6 with a composition Ti-5.8Al-4.0Sn-4.0Zr-0.7Nb-1.5Ta-0.4Si-0.06C was also tested [33]. Finally, finite element analysis was conducted by using ABAQUS 6.14 software to reveal the microscopic deformation mecha nism of the as-prepared alloys. The Mises stress, which is usually used to determine if a given material will yield or fracture, is calculated by using the equation: σ ¼ ð1=2ððσ1 σ 2 Þ2 þ ðσ2 σ3 Þ2 þ ðσ 3 σ 1 Þ2 ÞÞ1=2 , where σ 1, σ2 and σ3 is principal stress under different direction. The Young’s modulus of β-Ti phases and (Cu, Ni)Ti2 phase is 121 GPa and 42 GPa, respectively. The Poisson’s ratios for both phases were set as 0.3. These values were determined by nanoindentation experiment. The formation of microcrack was not considered in numerical simulation process. 3. Results
Fig. 1. XRD patterns of the Ti66Nb13Cu8Ni6.8Al6.2 sintered by using different powders and after deformation at elevated temperatures: (a) crystalline pow ders; (b) crystalline powder and after deformation at 500 � C; (c) amorphous powders; (b) amorphous powders and after deformation at 500 � C.
Fig. 1 displays the XRD pattern of Ti66Nb13Cu8Ni6.8Al6.2 alloys sin tered by using amorphous and crystalline powder, and after deformation 2
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Table 1 Compressive testing properties for the as-fabricated alloys at ambient and at elevated temperatures. The yield stress σy, ultimate stress σf, fracture strain εf. Herein “Alloy E” and “Alloy W00 denote the alloys with equiaxed and with Widmanst€ atten structures, respectively. Sample
T (oC)
σ y (MPa)
σf (MPa)
εf (%)
Alloy W
20 400 500 600
1380 � 15 1365 � 12 1260 � 16 610 � 10
1885 � 11 1500 � 16 1153 � 13 300 � 12
13.3 � 0.2 24.1 � 0.1 23.2 � 0.3 >50
Alloy E
20 400 500 600
1450 � 12 1360 � 8 853 � 15 400 � 14
2350 � 11 1405 � 13 869 � 13 239 � 12
28.5 � 0.2 35.2 � 0.1 >50 >50
TG6
400 500 600
950 � 9 697 � 13 650 � 12
905 � 24 700 � 12 550 � 22
20.4 � 0.1 27.5 � 0.3 >50
equiaxed structure are also larger than other advanced titanium alloys fabricated by melt solidification, including bulk metallic glass [37–39] and nanostructure-dendrite bimodal composites [4–8,16–18]. In €tten structure matrix exhibits higher contrast, the alloy with Widmansta yield strength and ultimate strength at elevated temperatures than those of the alloy with equiaxed bcc β-Ti matrix (Fig. 3c–d and Table 1). The yield strength, fracture strength and fracture strain of the alloy with €tten microstructure matrix at 500 � C is 1260 MPa, 1153 MPa, Widmansta and 23.2%, respectively (Fig. 3c), which are superior to those of the commercial Ti-alloy TG6. Besides, it is found that the yield strength of the alloy W decreases drastically as the testing temperature is increased to 600 � C, while this softening temperature point decreases to about 500 � C for the alloy E (Table 1). The results indicate the alloy W with €tten structure matrix have a better high temperature Widmansta resistance. In order to clarify origins of different mechanical properties for the as-fabricated alloys at room and elevated temperatures, TEM in vestigations were conducted on the specimens deformed (8%-strained) under different temperatures. As shown in Fig. 4a, nano-sized cracks arise at the grain boundaries between the (Cu, Ni)Ti2 phases and be tween the (Cu, Ni)Ti2 and β-Ti matrix in the alloy E. Meanwhile, abundant dislocations were clearly observed in the β-Ti matrix and hindered by the (Cu, Ni)Ti2 phases, but seldom were found in (Cu,Ni)Ti2 phases. The results indicate that the β-Ti phase regions have the excel lent plastic deformation and work-hardening capacity thereby acts as the ductile phases with respect to the (Cu, Ni)Ti2 phases. Besides, TEM observations show that there are lots of deformation-induced nanotwins (or stacking faults) in the equiaxed β-Ti matrix in alloy E (Fig. 4b). The width of the nanostructure determined by HRTEM is about eight inter atomic distances. The corresponding SAED pattern of the nanostructure exhibits elongated diffraction spots along the crystal planes (110) and (1ð Þ10) (Fig. 4b inset), indicating that there are severe lattice dis tortions inside this region. In contrast, the 8%-strained at 500 � C and then quenched sample exhibits completely different morphologies from the corresponding counterparts deformed at room temperature (Fig. 4c). Firstly, the equiaxed fcc (Cu, Ni)Ti2 phases in alloy E undergoes distinct plastic deformation at high temperature. This can be testified by their elliptic or irregular shapes as indicated by the red dashed lines (Fig. 4c). Secondly, no micro-cracks and damaged interface between the (Cu, Ni) Ti2 and β-Ti phases were observed in the high-temperature deformed sample (Fig. 4d), different from that formation of nano-sized cracks in the room-temperature deformed sample (Fig. 4a). The results indicates that the plastic deformation ability of (Cu, Ni)Ti2 phases increases drastically with increasing of testing temperature. The microstructures for the deformed Ti66Nb13Cu8Ni6.8Al6.2 com posite with Widmanst€ atten structure matrix are displayed in Fig. 5. Similar to the sample with equiaxed bcc β-Ti matrix, deformationinduced dislocations and cracking were formed in the β-Ti phases and
Fig. 2. SEM microstructures of the Ti66Nb13Cu8Ni6.8Al6.2 bulk alloys fabricated by sintering from (a) amorphous and (b) crystalline powder. (c) TEM image and (d) corresponding magnified Widmanst€ atten structure matrix of the alloy in (b).
crystallizing amorphous powders always displays the equiaxed struc ture, while the alloy fabricated by sintering of crystalline powder pos €tten structure. sesses a Widmansta Fig. 3 presents the compressive engineering stress-strain curves of the sintered Ti66Nb13Cu8Ni6.8Al6.2 bulk alloys with equiaxed bcc β-Ti €tten matrix (alloy W) at room matrix (alloy E) and with Widmansta temperature and at elevated temperature. For comparison, the stressstrain curves of commercial high-temperature titanium alloy TG6 is also displayed in Fig. 3. The corresponding mechanical properties are listed in Table 1. With increasing of test temperature, the yield strength and fracture strength decrease for both alloy E and alloy W (Table 1). Meanwhile, it is found that the alloy with equiaxed structure matrix has higher yield strength and ultimate strength and larger plasticity at room €tten structure ma temperature than those of the alloy with Widmansta trix (Fig. 3a–b and Table 1). The strength and plasticity in the alloy with
Fig. 3. Compressive engineering stress-strain curves at room temperature (a and b) and elevated temperature (c and d) of the fabricated Ti66Nb13Cu8 Ni6.8Al6.2 bulk alloys with different structures. Herein “E” and “W” denote the alloys with equiaxed and with Widmanst€ atten structure matrix, respectively. (e) The commercial high-temperature alloy TG6. 3
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plasticity in room-temperature. When the samples is deformed at 500 � C, interestingly, it can be found that the acicular (Cu, Ni)Ti2 phases are curved into the “C” or “S” shapes (Fig. 5c–d). Meanwhile, profuse dislocations are hindered and piled up in the β-Ti matrix where is close to the (Cu, Ni)Ti2 phases, showing that dislocations are hard to cross the boundary of β-Ti phase and (Cu, Ni)Ti2 phases. The corresponding SAED pattern confirms that “C” or “S” shaped phases possess fcc structure and the high temperature does not cause phase transformation of the (Cu,Ni) Ti2 phase (Fig. 5d), which is consistent with the XRD results (Fig. 1). Recently, finite element analysis (FMA) as a supplementary method was applied to reveal the essential relationship of microstructure and property of metallic materials. For instance, Jeon et al. employed the FMA method to simulate deformation behaviors of Ti-based amorphous matrix composites [40]. The FMA results of deformation bands or shear bands are in good agreement with the experimental findings [40]. In order to investigate further the effects of microstructure on the ductility of the present Ti66Nb13Cu8Ni6.8Al6.2 composites, the FMA was per formed to reveal the microscopic deformation mechanism. According to the above SEM and TEM observations, it can be concluded that the €tten structure matrix and equi biggest difference between Widmansta axed structure matrix is that there are acicular (Cu,Ni)Ti2 phases in the €tten structure but not in equiaxed structure. Therefore, the Widmansta numerical simulation models for the alloys with different microstructure are simplified as Fig. 6a and Fig. 6c, respectively. The white regions denote (Cu,Ni)Ti2 phases and blue regions donate the β-Ti phase. The Mises stress fields in the composites under different strains are displayed in Fig. 6b–c and Fig. 6e-f, respectively. When the strain is 0.6%, the maximum stress in the alloy with equiaxed β-Ti phase matrix is mainly distributed at junction of the (Cu,Ni)Ti2 phases and (Cu,Ni)Ti2 phases, or at interface between the (Cu,Ni)Ti2 phase and the matrix (Fig. 6b). The stress concentration increases with increase of strain (Fig. 6c). In the samples having acicular (Cu,Ni)Ti2 phases, the maximum stress also distributes at the boundary of the (Cu,Ni)Ti2 phases when the strain is small (Fig. 6e). However, the stress begins to concentrate in the middle of acicular (Cu,Ni)Ti2 when the sample is yielded (Fig. 6f). The nu merical simulation results indicate that acicular (Cu,Ni)Ti2 phases in the alloys play a critical key role on the stress distribution during deformation. The evolutions of Mises stress values extracted from the different positions in the samples are displayed in Fig. 7. During the elastic deformation stage, the stress values in the β-Ti phases are higher than those both in equiaxed and in acicular (Cu,Ni)Ti2 phases (Fig. 7a–b), which is mainly ascribed to the fact that elastic modulus of β-Ti phases is larger than that of (Cu,Ni)Ti2 phases [36]. With increase of loading, in alloy E with equiaxed structure the Mises stress at the interface between the (Cu,Ni)Ti2 phase and β-phase matrix increase sharply (Fig. 7a). When the strain is 8%, the maximum stress at the junction reaches about 2023 MPa, which is about 1.6 times of loading (Fig. 7a). The result in dicates that cracking prefers to form at interface between the (Cu,Ni)Ti2 phase and β-phase matrix in the alloy E with equiaxed structure, which agrees well with the TEM observations (Fig. 4a). In the sample W with Widmanst€ atten structure, Mises stress at interface between the (Cu,Ni) Ti2 phase and β-phase matrix is also larger than that in matrix or second phases. However, the Mises stress values in the middle of acicular (Cu, Ni)Ti2 phases increases rapidly (Fig. 7b). When the strain is 8%, the maximum values in the acicular (Cu,Ni)Ti2 phases can reach 3701 MPa, which is far larger than that at the interface of (Cu,Ni)Ti2 phases and β-phase matrix. The result indicates the cracks in the samples with Widmanstδtten structure will be firstly formed in the middle position of acicular (Cu,Ni)Ti2 phases.
Fig. 4. (a) TEM bright-field micrograph of the 8%-strained sample E with equiaxed bcc β-Ti matrix at room temperature, (b) Deformation induced nanotwins in β-Ti matrix. The insets are corresponding SAED pattern and highresolution TEM (HRTEM) micrograph of the square region respectively. (c) TEM bright-field micrograph of the 8%-strained sample E with equiaxed bcc β-Ti matrix at 500 � C, (d) High-resolution TEM image of the phase boundary of the β-Ti matrix and fcc (Cu, Ni)Ti2 phase in (c).
at the boundaries between the equiaxed second phase and matrix (Fig. 5a), respectively. However, a large number of micro-cracks was formed in the interior of acicular (Cu,Ni)Ti2 phases (Fig. 5b). The results indicate that acicular (Cu,Ni)Ti2 phases may be the origin of the low
4. Discussion
Fig. 5. (a) TEM bright-field micrograph of the 8%-strained sample with Wid manst€ atten structure matrix at room temperature, (b) micro-cracks in middle of acicular (Cu, Ni)Ti2 phase. (c–d) TEM morphologies of the acicular (Cu, Ni)Ti2 phases in the 8%-strained sample at 500 � C. The inset in (d) is the corre sponding SAED pattern of fcc (Cu, Ni)Ti2 phase.
4.1. Deformation at room temperature In the previous works, lots of equiaxed structural Ti–Nb/Ta–Cu–Ni/ 4
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Fig. 6. Finite element analysis models used for calculation of Ti-based composites with equiaxed structure (a) and Widmanst€ atten structure (d); Stress field of the equiaxed structural composites applied 0.6% strain (b) and 8% strain (c); Stress field of the Widmanst€ atten structural composites after being applied to 0.6% strain (e) and 8% strain (f).
Co–Al composites have been prepared successfully by crystallizing amorphous powder [20–27], but the origins of high strength and large plasticity are ambiguous. Although the microstructures including pile-up of dislocations and cracks at the interface between the TiFe phase and the matrix have been observed in deformed samples by Li et al. [41,42], up to now the detailed microscopic deformation mecha nism is not elaborated clearly, especially in the samples having acicular phases. In this work, the evolutions of Mises stress during deformation are revealed by FMA method (Fig. 7). By coupling with TEM results, the deformation mechanism of the present Ti66Nb13Cu8Ni6.8Al6.2 compos ites is proposed. The corresponding schematic diagram is displayed in Fig. 8. When loading is small, in the sample with equiaxed bcc β-Ti matrix, elastic deformation occurs simultaneously in both the β-Ti matrix and the (Cu,Ni)Ti2 phases, but the β-Ti phases present considerably higher stress than the (Cu, Ni)Ti2 phases due to the lower modulus of (Cu, Ni) Ti2 phases (Fig. 8a, corresponding to I stage in Fig. 7a). With increasing of stress beyond the yield point of the alloy, a large number of disloca tions are generated in the ductile β-Ti matrix and are pinned by the surrounded equiaxed (Cu, Ni)Ti2 phases (Fig. 8b, corresponding to II stage in Fig. 7a). The profuse dislocation pile-ups would largely contribute to the increase in strength due to working hardening effect [18]. With increasing further the loading, stress concentration increases drastically at the boundaries of equiaxed (Cu, Ni)Ti2 phases (Fig. 7a). When the stress level is increased to a critical value, cracks as shown in Fig. 4a would be generated at boundaries of the equiaxed (Cu, Ni)Ti2 phase regions (Fig. 8c, corresponding to the III stage in Fig. 7a). Since the (Cu, Ni)Ti2 phase in the Ti66Nb13Cu8Ni6.8Al6.2 alloys represents considerably lower flow stress and surrounded by β-Ti phases, the deformation of the (Cu, Ni)Ti2 phase is constrained by the β-Ti phase, consequently, the formed cracks at the boundary are separated and restricted by the adjacent ductile β-Ti matrix (Fig. 4a). Therefore, they are unable to shear-off the whole sample [29]. With increase of stress, the sustaining retention of the increasing number in dislocation pile-ups and strain induces formation of nanotwins in the ductile β-Ti matrix, as
shown in Fig. 4b. It was reported that twinning after an initial dislocation-meditated hardening process could allow the bcc β-Ti phase to undergo superplastic-like deformation [42], which would promote the increases of the plastic strain of samples (Fig. 8d, corresponding to the IV stage in Fig. 7a). When the load is increased to a high enough value that β-Ti matrix cannot suppress propagation of micro-cracks, the cracks would pass through the whole matrix, thus leading to the macroscopic fracture of the as-fabricated bulk alloy. Therefore, the equiaxed Ti66Nb13Cu8Ni6.8Al6.2 composites prepared by crystallizing amorphous powders have high fracture strength and large plasticity, and the excellent mechanical properties mainly originate from the special microstructure and corresponding deformation mechanism. Similar microscopic deformation process arises in the alloy with €tten structure matrix. The only difference is deformation Widmansta behavior of acicular (Cu, Ni)Ti2 phases. It was recognized that hardness of β-Ti phase was three times that of (Cu, Ni)Ti2 phase [36]. This means that the strength of (Cu, Ni)Ti2 phase is lower than that of β-Ti phase, and thus (Cu, Ni)Ti2 phase is easier to break than β-Ti phase. Therefore, €tten structure can the deformation process for the alloy with Widmansta be divided into following several stages. When the loading is small, the stress concentration distributes in the interface between the (Cu,Ni)Ti2 phase and β-phase matrix (Fig. 8a, corresponding to I stage in Fig. 7b). With increasing of loading, stress concentration increases drastically at the interior of acicular (Cu, Ni)Ti2 phases. Subsequently, the cracks would be formed in the acicular (Cu, Ni)Ti2 phases (Fig. 8b, corre sponding to II stage in Fig. 7b). As the stress increases further, larger sized cracks will be formed at the boundary of micron-sized (Cu, Ni)Ti2 phases like that in alloy with equiaxed structure (Fig. 8c, corresponding to III stage in Fig. 7b). Since the density of acicular (Cu, Ni)Ti2 phases in €tten structure matrix is high, the formed cracks in the acic Widmansta ular phase would deteriorate plastic deformation ability of Wid manst€ atten structure matrix. Therefore, once the crack is formed in boundary of micron-sized (Cu, Ni)Ti2 phases, they are hard to be sup pressed by the adjacent Widmanst€ atten structure matrix having micro cracks. As a result, the samples will break quickly with low plasticity and 5
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strength (Fig. 8d, corresponding to IV stage in Fig. 7b). Therefore, the €tten structure matrix ex Ti66Nb13Cu8Ni6.8Al6.2 alloy with Widmansta hibits lower strength and less plasticity than that with equiaxed bcc β-Ti matrix at room temperature. 4.2. Deformation at elevated temperature As shown in Fig. 3 and Table 1, the yield strength of the two samples with equiaxed bcc β-Ti matrix and with Widmanst€ atten structure matrix decreases with increasing the testing temperature. The yield point is always a critical stress to start emission of dislocation or activation of sliding. It is widely recognized that the critical stress decreases with increasing the temperature [43,44]. Besides, the yield strength in the composite is also determined by second phases [45]. TEM observation shows that the equiaxed (Cu, Ni)Ti2 second phases are compressed into elliptic or irregular shapes under deformation, indicating that strength of the (Cu, Ni)Ti2 second phases decreases sharply at the high temper ature. Therefore, the yield strength for each sample decreases with increasing the testing temperature. With regard to the plasticity, it is associated with the deformation mechanism. It has been reported that the glide of grains or phase boundaries becomes easier with increasing temperature [43,44]. TEM observations also demonstrate that the deformation abilities of the (Cu, Ni)Ti2 and β-Ti phases increase with increasing the temperature (Fig. 4c–d and Fig. 5c–d). The increase of plastic deformation capacity will lead to release of the stress concen tration resulting in crack nucleation in grain boundaries. Therefore, the plasticity increased with the increasing test temperature for each specimen. Also, it can be noticed that the sample with Widmanst€ atten structure matrix exhibits higher yield strength and fracture strength than the alloy ^-Ti matrix at elevated temperatures, indicating its with equiaxed bcc a higher deformation resistance ability to elevated temperatures. This is ^-Ti phase content in alloys with Widmansta €tten mainly ascribed to the a structure and deformation mechanism of the acicular (Cu,Ni)Ti2 phases. Under the high temperature, it was reported that the minimum sliding resistance in titanium alloy is the β phase/β phase boundary [46,47], and only the dislocations which slipped along β-Ti slip system can cross the β/α interfaces easily [33]. As shown in Fig. 2d, the β-Ti phases in €tten structure are interleaved by the acicular (Cu, Ni)Ti2 Widmansta phases. Besides, the volume fraction of β-Ti in the alloys with Wid manst€ atten structure matrix is far lower than that in the samples with equiaxed bcc β-Ti matrix. Therefore, the content of β/β boundary in the €tten structure is less than that in alloy with equi alloy with Widmansta axed structure. So the alloy with Widmanst€ atten structure matrix pos sesses higher yield strength at elevated temperatures (Table 1). Also, it can be seen that the equiaxed (Cu, Ni)Ti2 phases are elongated at high temperature (Fig. 4c), while the curving of the acicular (Cu, Ni)Ti2 phases is observed in Widmanst€ atten structure matrix (Fig. 5c–d). This indicates that the acicular (Cu, Ni)Ti2 phases can hinder the €tten higher-temperature flow behavior of the sample with Widmansta matrix thereby result in a higher ultimate strength. Shortly, the presence €tten structure matrix of the acicular (Cu, Ni)Ti2 phases inside Widmansta is the decisive reason for the higher yield and ultimate strength at elevated temperature.
Fig. 7. Mises stress extracted from the numerical simulation results of alloys with different structure: (a) equiaxed structure; (b) Widmanst€ atten structure. The “1”, “2”, “3”, and “4” denotes different positions marked in Fig. 5a and d.
5. Conclusion Deformation mechanism of Ti66Nb13Cu8Ni6.8Al6.2 composites with equiaxed β-Ti and with Widmanst€ atten structure matrix was investi gated at ambient and elevated temperatures. Results show that the equiaxed β-Ti structure matrix has higher capability of plastic defor mation with formed nanotwins, while cracking arises in the interior of acicular (Cu, Ni)Ti2 phases in Widmanst€ atten structure matrix. This leads to the higher strength and plasticity of the sample with equiaxed structure matrix at ambient temperature. In contrast, abundant dislo cations generated in β-Ti phases are hindered by curved acicular (Cu, Ni)
Fig. 8. Schematic diagram of deformation behavior in Ti66Nb13Cu8Ni6.8Al6.2 bulk alloys with two different matrixes at room temperature. The left and right ones in each figure denote the alloy with Widmanst€ atten structure matrix and with equiaxed bcc β-Ti matrix respectively. The “I”, “II”, “III” and “IV” stage correspond to the different regions in Fig. 6. 6
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€tten structure matrix at elevated tempera Ti2 phases inside Widmansta €tten tures. This causes the higher strength of the sample with Widmansta structure matrix. The results obtained would provide some insights into designing high-performance metallic alloys used as high-temperature structural materials.
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