- Email: [email protected]
Ti composites dominated by primary structure of Ti matrix
Ceramics International 45 (2019) 17767–17774
Contents lists available at ScienceDirect
Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Interfacial reactions and matrix microstructure evolution in SiCf/Ti composites dominated by primary structure of Ti matrix
T
Shuming Zhanga, Minjuan Wangb, Mao Wenc, Ming Wuc, Qingfeng Wanga,∗, Hao Huangb,∗∗ a
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, 066004, China AECC Beijing Institute of Aeronautical Materials, Beijing, 81-15, 100095, China c State Key Laboratory of Superhard Materials, Department of Materials Science, Key Laboratory of Automobile Materials, MOE, Jilin University, Changchun, 130012, China b
ARTICLE INFO
ABSTRACT
Keywords: SiCf/Ti composite Microstructure Interfacial reaction Strength
High-performance continuous SiC fiber reinforced titanium matrix composites (SiCf/Ti) are strongly required by aerospace vehicles, and their properties strongly depend on the microstructure of the interfacial reaction layer and the matrix. Thus, it is important to explore new routes to regulate these two SiCf/Ti components. In this work, α-phase-dominated and β-phase-dominated primary PVD Ti17 coatings were fabricated by controlling the bias voltage during physical vapor deposition (PVD). Their effects on the matrix microstructure evolution and interfacial reactions in SiCf/Ti after consolidation were further explored. It was found that compared with a primary α-Ti PVD Ti17 coating, a primary β-Ti PVD coating induced competitive growth between α-Ti and β-Ti grains in SiCf/Ti17 during consolidation. This yielded a smaller grain size and more β phase in the SiCf/Ti17 evolved from the primary β-Ti coating. In addition, a thicker reaction layer formed, as well as a lower interface strength and tensile strength in the SiCf/Ti17 evolved from a primary β-Ti coating, compared with the SiCf/Ti17 evolved from a primary α-Ti PVD Ti17 coating.
1. Introduction Light-weight high-strength structural materials, especially those possessing high specific strength and stiffness at both room and elevated temperatures, are highly desired for the construction of aerospace vehicles and advanced propulsion systems. Continuous SiC fiber reinforced titanium matrix composites (SiCf/Ti) have been explored as potential candidates in such applications [1–3]. It is generally accepted that interfaces play a crucial role in determining the properties of SiCf/ Ti, such as tensile strength and fatigue failure [4,5]. Thus, many researchers have investigated and tailored interfacial behaviors of SiCf/Ti to improve their performance via barrier layer optimization and adjustment of consolidation processes [6–9]. Wu et al. improved the tensile strength of SiCf/Ti through optimization of its interfacial properties by controlling the texture of the carbon coating, in which the distribution of interfacial reaction zones evolved from continuous elemental diffusion and grain growth during consolidation process was governed by the texture of the C coatings, consisting of turbostratic amorphous carbon, fine-grained carbon, transition TiC, coarse-grained TiC, and TiC with a decreasing C content gradient [10]. In addition, by
∗
adjusting the hot isostatic pressing (HIP)consolidation temperature, Wang et al. observed that the reduce or disappearance of fine TiCx layer thickness but rapid increment of the thickness of total reaction layers with elevating the consolidation temperature [11]. Barrier coatings and consolidation processes are commonly applied to control the thickness and microstructure of interfacial reactions, allowing the mechanical properties of SiCf/Ti composites to be further tailored. Since the interfacial microstructure strongly influences the properties of SiCf/Ti [12,13], there is a need to explore how interfacial reactions are affected by other avenues. Notably, during the deformation and failure of SiCf/Ti, the matrix usually provides nucleation sites for dislocations and cracks, and the subsequent expansion to interface zones can induce interfacial debonding [14,15]. This means that in addition to the interfacial microstructure, the matrix may also play an important role in influencing the mechanical properties of SiCf/Ti. Some recent attempts have been made to regulate the matrix microstructure by changing the consolidation temperature and duration [16–18]. The matrix-coated fiber (MCF) method has been widely utilized to prepare high-performance SiCf/Ti over the past several few years due to its unique advantages [19,20], in
Corresponding author. Corresponding author. E-mail addresses: [email protected] (Q. Wang), [email protected] (H. Huang).
∗∗
https://doi.org/10.1016/j.ceramint.2019.05.347 Received 13 March 2019; Received in revised form 27 May 2019; Accepted 30 May 2019 Available online 31 May 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
which a Ti alloy coating is deposited onto SiC fibers by physical vapor deposition (PVD) to serve as the primary structure of matrix. Obviously, PVD Ti alloy, directly contacted with SiC fibers, would become as both the reaction front and matrix source whose microstructure would strongly influence the interfacial reaction and matrix microstructure. However, the interfacial reaction and matrix microstructure in the SiCf/ Ti dominated by primary microstructure of PVD Ti alloy is scarcely investigated. In this paper, α-phase-dominated and β-phase-dominated primary PVD Ti17 coatings were prepared by controlling the bias voltage during PVD, and their effects on the interfacial reaction and matrix microstructure evolution in the SiCf/Ti after consolidation were also explored. The results showed that the primary structure of the PVD Ti17 coating determined both the interfacial reaction and matrix microstructure of SiCf/Ti. This work can provide valuable insight into the influence of the PVD Ti alloy structure on the matrix microstructure and interfacial reaction to further improve the performance of the SiCf/Ti. 2. Experimental procedures and characterization 2.1. Sample preparation The MCF method was used to fabricate Ti17 coatings onto tungstencored SiC monofilaments as precursor wires in a facing-targets magnetron-sputtering system, in which the Ti17 alloy (nominal composition in wt.%: 4.5–5.5 Al, 1.6–2.4 Sn, 1.6–2.4 Zr, 3.5–4.5 Mo, 3.5–4.5 Cr, and balance Ti) was utilized as the target, and C-coated SiC fibers with ∼100 μm diameters were used as the substrate. By regulating bias voltage, both α-phase- and β-phase-dominated PVD Ti coatings were successfully fabricated, and used as the primary matrix structure to be further evolved into an SiCf/Ti matrix after consolidating. Next, these precursor wires were placed into a Ti17 alloy canister hermetically to conduct sealing process using electron beam wielding. Subsequently, SiCf/Ti17 was manufactured by consolidating these packed precursor wires which have been sealed in a canister through HIP under the condition of 920 °C/120MPa/120mins, and then the furnace was cooled to room temperature. After HIP, both α-phase- and β-phase-dominated SiCf/Ti17 precursor wires, named as SiCf/α-Ti17 and SiCf/β-Ti17 respectively, were intact and compacted as shown in Fig. 1, and exhibited highly uniform distributions of SiC fibers. 2.2. Characterization Tensile samples were prepared for each HIP consolidated specimen with gauge dimensions of Φ3 × M6. Room tensile tests were performed at an extension rate of 1 mm/min, using an Inspekt Table 100 kN model universal testing machine. PVD titanium alloy was sputtered on a single-crystal silicon wafer, and one piece was peeled off for microstructure characterization. Primary Ti matrices were observed in parallel with the direction of columnar crystal growth. Both X-ray diffraction (XRD, D/max-2500) and field emission JEOL 2010F transmission electron microscopy (TEM) were employed to characterize
PVD Ti17 coatings to identify the primary phase structure of the matrix. The TEM specimens were thinned by a Leica EM RES 102 ion beam milling system. The interfacial reaction products were investigated by TEM, and the chemical compositions of each sub-layer in the interfacial zones were further examined by energy dispersive X-ray spectroscopy (EDS) equipped to the TEM. TEM specimens were cut from the composites perpendicular to the cross-section and in the radial direction of the fiber by FEI Quanta 200 FEG focused ion beam (FIB) milling in the process shown in Fig. 2. Scanning electron microscopy (SEM, FEI nano 450) was utilized to obtain micrographs of the samples. SiCf/Ti17 specimens used for SEM observations were cut along the cross-section, polished, and etched in an HF solution. The interfacial debonding strength was evaluated by performing fiber push-out tests on the sample slices with 0.5 ± 0.01 mm thicknesses [21]. Fractographies were observed by SEM and laser scanning confocal microscopy (LSCM, KEYENCE VK-100). 3. Results and discussion The nanocrystalline nature of PVD Ti alloys, containing large numbers of grain boundaries and different phase structures, would significantly affect the interfacial reactions and matrix microstructure of SiCf/Ti during consolidation. Two typical phase structures, including α-Ti nanocrystalline and β-Ti nanocrystalline, can be first obtained in the primary PVD Ti coating by regulating the PVD conditions to investigate its effects on the interfacial reactions and matrix evolution of SiCf/Ti. 3.1. Microstructure of primary PVD Ti17 coating As expected, α-Ti and β-Ti structures were obtained in the primary PVD Ti17 coatings fabricated at −150 V and −250 V of bias voltage, respectively, as shown by XRD and selected area electron diffraction (SAED)results. The Ti17 coatings were first analyzed by XRD, and the resulting spetra are shown in Fig. 3. At −150 V, a strong peak located at 36.2° appears, corresponding to hcp α-Ti 101¯0 , a stable structure at room temperature. Conversely, at −250 V, the α-Ti 101¯0 peak disappeared, and a weak and broad peak centered at 37.6° appeared, which was attributed scribed to bcc β-Ti 110 . This implies that besides the crystal structure, the high bias voltage of −250 V also reduced the crystallinity or created smaller grain sizes relative to −150 V due to the higher incident energy, which induced re-nucleation. The average grain size was estimated from the full-width half maximum in Fig. 3 to be ∼15 nm for −150 V and ∼4 nm for −250 V. Moreover, due to the higher compressive stress introduced by the bombardment of higherenergy particles (−250 V), the metastable β-Ti nucleation and growth was stabilized [22]. The SAED pattern in Fig. 4 further suggested that an α-Ti structure formed at −150 V and a β-Ti structure formed at −250 V, by assigning all patterns in −150 V and −250 V to α-Ti and βTi, respectively. The HRTEM images in Fig. 4 show typical polycrystalline character for both coatings, in which the grain size is 10–20 nm for −150 V and 3–8 nm for −250 V accompanied by partial amorphous zones, in agreement with the XRD results. The cross-sectional morphologies of the precursor wires were also investigated by SEM (Fig. 4b and d). Low-energy particle bombardment at −150 V resulted in typical column growth due to the shadow effect, whereas higher-energy energy particle bombardment at −250 V promoted the formation of dense structures due to higher diffusion and ion sub-implantation effects. 3.2. Microstructure of SiCf/Ti17
Fig. 1. Cross-sectional SEM micrographs of (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17.
3.2.1. Microstructure of titanium alloy matrix After being consolidated at the same HIP condition (in the α+β phase fields), the fabricated α-Ti and β-Ti Ti17 coatings on SiC fibers mainly evolved into corresponding SiCf/α-Ti17 and SiCf/β-Ti17
17768
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
Fig. 2. The process used to cut SiCf/α-Ti17 samples by FIB for (a)–(d).
Fig. 3. XRD patterns of primary PVD Ti17 coatings.
matrices in addition to forming some interfacial zones. The effect of the primary phase structure of PVD Ti on the interfacial reaction and matrix microstructure of SiCf/α-Ti17 and SiCf/β-Ti17 is clarified as follows. The representative matrix microstructures of SiCf/α-Ti17 and SiCf/βTi17 respectively evolved from primary α-Ti and β-Ti observed by SEM, as shown in Fig. 5. It was found that both composites exhibited typical equiaxed structures, but there were large differences in the grain sizes and content of each phase (α phase and β phase). The grain size of equiaxed α phase in the SiCf/α-Ti17 was much larger than in SiCf/βTi17, and SiCf/α-Ti17 also contained less β phase than SiCf/β-Ti17 (34% β phase for SiCf/α-Ti17 and 43% for SiCf/β-Ti17). The above differences in grain size and phase content in the SiCf/α-Ti17 and SiCf/ β-Ti17 mainly arose from the different growth behaviors during consolidation. Notably, as the α+β phase fields were applied to HIP
consolidation, both continuous grain growth and temperature-induced phase transition of matrix simultaneously occurred during the heating and isothermal stages. During the heating stage of SiCf/β-Ti17 consolidation, the primary metastable β-Ti phase became unstable and partially transformed into stable α-Ti, and the newly formed α-Ti coexisted with the existing β-Ti to impede the grain growth of each phase. In contrast, during the consolidation of SiCf/α-Ti17, the primary stable α-Ti phase tended to continually grow without phase transition before reaching the temperature of α+β phase field. That is to say, the competitive growth between α-Ti and β-Ti grains in SiCf/β-Ti17 during heating gave rise to smaller grains and more β phase before entering into the isothermal stage, compared with SiCf/α-Ti17. At the isothermal stage (α+β phase field), the α phase began to partially transform into β phase, and each phase continued to grow. Therefore, compared with SiCf/α-Ti17, the remarkable competitive grain growth between α-Ti and β-Ti grains in SiCf/β-Ti17 during heating resulted in much finer equiaxed α-Ti grains. Meanwhile, residual primary β-Ti grains in SiCf/ β-Ti17 also contributed to more β phase. 3.2.2. Microstructure of interfacial reaction layer Besides the matrix structure of SiCf/Ti, the primary phase structure of PVD Ti17 also strongly influenced the evolution of interfacial reaction layer (RL) during HIP consolidation. From the SEM images in Fig. 5, it is observed that RL in both SiCf/α-Ti17 and SiCf/β-Ti17, located between the C coating and Ti17 matrix, were compact without any cracks or holes, but a thicker RL was observed in SiCf/β-Ti17 (∼1.7 μm for SiCf/α-Ti17 and ∼2.4 μm for SiCf/β-Ti17). It is accepted that the RL mainly arose from the continuous diffusion of elemental carbon from the carbon coating into the adjacent Ti matrix and to form TiCx products during high-temperature consolidation, known as Ti + C = TiC [23,24]. The distribution of reaction products in the RL was confirmed to play an important role on the interfacial properties
17769
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
Fig. 4. The HRTEM images and SAED patterns of (a) SiCf/α-Ti17 and (c) SiCf/β-Ti17; the cross-sectional images of Ti17 precursor wires of (b) SiCf/α-Ti17 and (d) SiCf/β-Ti17.
Fig. 5. Microstructure of SiCf/Ti17 observed by SEM for (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17.
[12,25]. To clarify the product distribution and PVD Ti structuredominated interfacial reactions, the interfacial reaction zones were further investigated by TEM. The TEM images are shown in Fig. 6. As determined using SAED patterns, only the TiCx phase that was found in the RL of both composites was located between the C coating and the Ti17 matrix. However, the grain size of TiCx products showed a broad distribution, ranging from several nanometers near the C coating to hundreds of nanometers or more close to the Ti17 matrix, nearly approaching the grain size of the Ti17 matrix. Based on the TiC grain sizes, the RL can be divided into three sub-layers, namely Ⅰ RL, Ⅱ RL, and Ⅲ RL, which respectively correspond to fine-grained layers adjacent to the C coating, the transition layer in the middle of the RL zone, and the coarse-grained layer next to the matrix [10]. The corresponding thicknesses and grain sizes of the three sub-layers in the RL region and grain size of the Ti matrix for the two samples are listed in Table 1. For SiCf/α-Ti17, the total thickness of RL was determined to be∼1.7 μm,
consisting of ∼60 nm Ⅰ RL, ∼200 nm Ⅱ RL, and∼1.4 μm Ⅲ RL. Meanwhile, the grain size rapidly increased from less than 20 nm in Ⅰ RL, to ∼200 nm in Ⅱ RL, and larger than 500 nm in Ⅲ RL. The rapid increase in grain size along RL can be explained by noting that C atoms continuously diffused from the C coating towards the matrix, whose grains also grew during HIP. It has been previously confirmed that amorphous C usually exists in Ⅰ RL near the C coating due to the rich C source [26,27]. It impedes the TiC grain growth and maintains fine TiC grains (several nanometers), supported by the EDS result that C is abundant in Ⅰ RL (Table 2). It is clear that Ⅰ RL forms during the initial HIP stage, in which the C-rich character and nanocrystalline-boundaries can provide short-circuit paths for the diffusion of C atoms [28], which is described as a reaction-controlled process [1]. During the formation of Ⅰ RL, grain growth simultaneously occurred at zones near Ⅰ RL. As the reaction continues, the entry of C atoms across Ⅰ RL yields Ⅱ RL, which is responsible for the abrupt increase in grain size to ∼200 nm in Ⅱ RL. The newly formed Ⅱ RL with a much larger grain size exhibits nearstoichiometric character, as shown by the EDS result, and it can serve as a diffusion barrier layer [29]. Since the diffusion rate of C in the TiC lattice is a crucial route for C migration across Ⅱ RL, further interfacial reactions would be diffusion-controlled processes [1]. Obviously, the diffusivity rate of C in the TiC lattice is several orders of magnitude lower than in the Ti lattice [24] and is a short-circuit process. After C atoms diffuse across the Ⅰ RL and Ⅱ RL via further reactions, they would react with the grown Ti grains to form coarse-grained Ⅲ RL with grain sizes larger than 500 nm. The same variety of grain sizes along the RL also appeared in SiCf/βTi17, increasing from less than 20 nm in Ⅰ RL to larger than 200 nm in Ⅲ RL due to the combined roles of continuous C diffusion towards the matrix and the simultaneous grain growth of the matrix. The thickness of each sub-layer in SiCf/β-Ti17 was evaluated to be approximately 100 nm for Ⅰ RL, 400 nm for Ⅱ RL, and 1.9 μm for Ⅲ RL, which were ∼2.4 μm thicker than the RLs in SiCf/α-Ti17 (1.7 μm). Moreover, the grain size of each sublayer in SiCf/β-Ti17 was also much smaller than SiCf/α-Ti17, meaning that the primary α-Ti PVD precursor resulted in much larger grain sizes but thinner RLs compared with the primary β-Ti PVD precursor. Since the β-Ti PVD precursor was used as a reaction front, the additional phase transition from β-Ti to stable α-Ti occurred during interfacial reactions, yielding smaller matrix grains due to
17770
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
Table 2 EDS results of three sub-layers in the RL region of SiCf/α-Ti17 and SiCf/β-Ti17. SiCf/α-Ti17
Spot 1 (RL I) Spot 2 (RL II) Spot 3 (RL III)
SiCf/β-Ti17
C (at.%)
Ti (at.%)
C (at.%)
Ti (at.%)
64.27 42.13 28.05
28.36 55.98 67.15
59.77 38.36 28.93
34.86 61.22 69.67
RL- Reaction Layer.
Fig. 6. TEM brightfield images and SAED patterns of interfacial reaction products and Ti17 matrix for (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17.
competitive grain growth. After C diffused into the nearby matrix to form a RL, the smaller grain sizes also appeared in the RL of SiCf/βTi17. In addition, more β-Ti phase was present in the SiCf/β-Ti17 matrix relative to SiCf/α-Ti17, as revealed by the SEM images in Fig. 5. It has been previously reported that β titanium alloy matrices react more rapidly with SiC fiber than α titanium alloys due to the faster diffusion rate of C or Si in a β-Ti matrix [24]. Thus, the faster C diffusion in β-Ti
than α-Ti would be mainly responsible for the formation of a thicker RL. The smaller grain size also partially contributed to a faster diffusion rate due to more terrible crystalline-boundary diffusion. In the two samples, the elemental distribution of carbon and titanium along the growth direction of the reaction layer from spot 1 to spot 3, respectively, are given in Table 2. The data shows that Ⅰ RL is carbon rich, near-stoichiometric in Ⅱ RL, and C-poor in Ⅲ RL. The carbon content of each sample decreased as the distance away from C coating increased, and conversely increased with titanium, further supporting the above viewpoint of interdiffusion processes. The STEM images and EDS elemental maps were obtained to further visualize the elemental distribution of the interfacial zones of both SiCf/ α-Ti17 and SiCf/β-Ti17, as shown in Fig. 7. The matrix elements were almost completely blocked outside the C coating, but C atoms obviously diffused from the C coating towards the matrix and dominated the formation of interfacial reaction zones with gradient C content. Additionally, the clear boundaries of carbon coating, interfacial reaction zones, and titanium alloy matrix were distinguished in both SiCf/α-Ti17 and SiCf/β-Ti17, suggesting the same RL thickness obtained from the SEM images in Fig. 5. During HIP, the formation of interfacial reaction zones is mainly controlled by the continuous diffusion of C atoms from the C coating into the adjacent Ti matrix. The content and grain size of each phase in the Ti matrix near the C diffusion front also determines the C diffusion rate. The microstructure evolution of the matrix in primary α-Ti or β-Ti nanocrystalines was mainly dominated by the continuous grain growth and temperature-induced phase transition, as shown in Fig. 8. The main difference in the SiCf/α-Ti17 and SiCf/β-Ti17 matrices were due to the heating stage T1-T2, in which competitive growth between α-Ti and βTi grains occurred in SiCf/β-Ti17, yielding smaller grain sizes and more β phase than SiCf/α-Ti17 before entering T3. Consequently, the smaller grain size and additional β phase in SiCf/β-Ti17 provided more terrible crystalline-boundary diffusion and faster C diffusion in β-Ti relative to SiCf/α-Ti17. This resulted in the formation of a thicker RL and fewer and finer equiaxed α-Ti grains in SiCf/β-Ti17, compared with SiCf/αTi17, as illustrated in Fig. 8. 3.3. Interfacial strength and tensile strength Since primary α-Ti and β-Ti PVD coatings were used to fabricate SiCf/α-Ti17 and SiCf/β-Ti17, different interfacial RLs, including different grain sizes and thicknesses of each sublayer were observed in each of the composites. This would further enhance the interface strength of SiCf/α-Ti17 and SiCf/β-Ti17. It has been widely accepted that the interface strength of SiCf/Ti is a crucial parameter governing
Table 1 The corresponding thicknesses and grain sizes of three sub-layers in the reaction layer region and grain sizes of the Ti matrix for two samples. Thickness of RL
SiCf/α-Ti17 SiCf/β-Ti17
Grain size of TiC
Grain size of Ti matrix/μm
RL I/nm
RL II/nm
RL III/μm
RL I/nm
RL II/nm
RL III/nm
60 100
200 400
1.4 1.9
≤20 ≤20
= Thickness of RL II 20 ≤ TiC≤ 100
≥500 ≥200
RL- Reaction Layer. 17771
0.4–2 0.1–0.4
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
Fig. 7. STEM images and EDS elemental maps of the interfacial zones of (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17.
the failure modes of composites under external loads, and therefore has a significant influence on mechanical properties [30]. Thus, push-out experiments were carried out to evaluate the interfacial properties of SiCf/α-Ti17 and SiCf/β-Ti17, as shown in Fig. 9. In the load-displacement curves, the peak load Pd is considered to be the point at which the load corresponds to the onset of fiber protrusion from the back surface, and the average Pd of SiCf/α-Ti17 and SiCf/β-Ti17 was measured to be ∼17.5 N and ∼13.1 N, respectively. By substituting Pd values, the shear stress or interface strength τd can be calculated based on the following formula [31]: d
= Pd/ df h
(1)
where df is the diameter of the SiC fiber and h is the thickness of the composite. The calculated τd was ∼111 MPa for SiCf/α-Ti17 and ∼82 MPa for SiCf/β-Ti17, which approach the reported values in other systems measured using the same method. Fu et al. reported a τd of ∼118 MPa for an SiCf/C/Ti–6Al–4V system [7] and D. Osborne et al.
Fig. 9. Typical fiber push-out load-displacement curves for SiCf/α-Ti17 and SiCf/β-Ti17.
Fig. 8. Models of titanium alloy crystal growth of different structure PVD precursors during heat treatment, showing how it affects the C diffusion and reaction interface formation.
17772
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
in the RL of SiCf/β-Ti17. It is clear that the SiCf/α-Ti17 with a thinner RL exhibited a higher interfacial strength than SiCf/β-Ti17, and when RL thickness of SiCf/β-Ti17 reached 2.4 μm it may have promoted preferential crack nucleation at the RL Tensile tests were performed on SiCf/α-Ti17 and SiCf/β-Ti17, and the resulting curves are shown in Fig. 11. The average tensile strengths of SiCf/α-Ti17 and SiCf/β-Ti17 were calculated to be 2094 MPa and 1910 MPa, respectively, and each composite had a much higher tensile strength than the Ti17 matrix (1060 MPa). In the composite, both the SiC fibers and the interfacial properties significantly affected the tensile strength. The tensile strength of SiCf/α-Ti17 was ∼200 MPa higher than SiCf/β-Ti17 despite having the same fiber volume fractions. To further analyze this result, the fracture morphologies of SiCf/α-Ti17 and SiCf/β-Ti17 were observed as shown in Fig. 12. Both composites show typical step-like fractures involving obvious push-out phenomena of their short fibers, rather than either brush-like fracture caused by a too-low τd or brittle facture induced by a too-high τd due to lower interfacial debonding. The morphologies suggested that each composite should possess suitable interface strength, as evidenced by their high tensile strengths (> 1900 MPa). More detailed observations showed that the lower τd of SiCf/β-Ti17 resulted in the appearance of obviously longer push-out fibers than SiCf/α-Ti17, which agreed with classical shear-lag model predictions. This may be responsible for the lower tensile strength of SiCf/β-Ti17 since it has a lower τd than SiCf/α-Ti17. Therefore, by only controlling the primary structure of PVD Ti17 coatings (α-Ti or β-Ti), the interfacial reactions and microstructure of SiCf/Ti17 can be tailored to achieve a higher τd and tensile strength in SiCf/α-Ti17 than SiCf/β-Ti17.
Fig. 10. SEM micrographs of interfacial cracks in SiCf/β-Ti17 for (a) and (b).
Fig. 11. Typical tensile stress–strain curves of SiCf/α-Ti17 and SiCf/β-Ti17.
4. Conclusions Both α-phase-dominated and β-phase-dominated primary PVD Ti17 coatings were successfully fabricated onto SiC fibers by only regulating the bias voltage, followed by consolidation into SiCf/α-Ti17 and SiCf/βTi17 via HIP. Based on the microstructure observations and mechanical properties of the two composites, the following conclusions can be drawn: 1) During HIP, competitive growth between α-Ti and β-Ti grains appeared in the SiCf/β-Ti17 during the heating stage because primary β-Ti grains became unstable, yielding smaller grain sizes and more β phase in SiCf/β-Ti17 than SiCf/α-Ti17. Compared with SiCf/β-Ti17, the smaller grain size and higher β phase in the matrix of SiCf/βTi17 also induced more terrible crystalline-boundary diffusion and faster C diffusion in β-Ti. This resulted in the formation of thicker RLs consisting of a fine-grained layer, a transition layer, and a coarse-grained layer. 2) SiCf/α-Ti17 with a thinner RL exhibited a higher interface strength than SiCf/β-Ti17, when the RL thickness of SiCf/β-Ti17 reached 2.4 μm which may have promoted preferential crack nucleation at the RL. The lower interfacial strength of SiCf/β-Ti17 caused increased fiber push-out than in SiCf/α-Ti17, contributed to the lower tensile strength in SiCf/β-Ti17. This provides a new strategy to regulate the matrix microstructure and RL of SiCf/Ti to further tailor their properties.
Fig. 12. Fracture morphologies of (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17 as observed by SEM; fracture morphologies of (c) SiCf/α-Ti17 and (d) SiCf/β-Ti17 observed by LSCM.
observed a slightly higher τd of ∼165 MPa in an SCS-6/Timetal-21S system [25]. Even if same test method was employed to characterize τd many factors such as thickness, polishing, and the material systems of the tested samples would affect these values. Thus, it is difficult to quantitatively determine the optimum value of τd for SiCf/Ti composite systems, but a moderate τd is strongly preferred to effectively transfer the load of the matrix to SiC fibers and also to facilitate interfacial debonding instead of across the fibers [30]. When τd is too low, a brushlike fracture mode occurs, while a too high τd induces typical flat brittle mode fracture [13]. It is generally believed that the τd of composites can be determined by the RL thickness. RL with appropriate thicknesses can improve the τd, but above a critical thickness, τd would be worsened due to preferential crack nucleation occurring at the brittle RL during loading [32]. Fig. 10 shows the appearance of interfacial cracks
Acknowledgements This work was supported by the National Key Research and Development Program of China (Grant No. 2016YFB0301203) and the National Natural Science Foundation of China (Grant No. 51871109). References
17773
[1] Z.X. Guo, B. Derby, Solid-state fabrication and interfaces of fibre reinforced metal matrix composites, Prog. Mater. Sci. 39 (1995) 411–495, https://doi.org/10.1016/
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al. 0079-6425(95)00002-X. [2] J.M. Larsen, S.M. Russ, J.W. Jones, An evaluation of fiber-reinforced titanium matrix composites for advanced high-temperature aerospace applications, Metall. Mater. Trans. A 26 (1995) 3211–3223, https://doi.org/10.1007/BF02669450. [3] C. Leyens, F. Kocian, J. Hausmann, W.A. Kaysser, Materials and design concepts for high performance compressor components, Aero. Sci. Technol. 7 (2003) 201–210, https://doi.org/10.1016/S1270-9638(02)00013-5. [4] Y.Q. Yang, H.J. Dudek, J. Kumpfert, Interfacial reaction and stability of SCS-6 SiC/ Ti–25Al–10Nb–3V–1Mo composites, Mater. Sci. Eng. A 246 (1998) 213–220, https://doi.org/10.1016/S0921-5093(97)00749-1. [5] G.H. Feng, Y.Q. Yang, X. Luo, J. Li, B. Huang, Y. Chen, Fatigue properties and fracture analysis of a SiC fiber-reinforced titanium matrix composite, Compos. B Eng. 68 (2015) 336–342, https://doi.org/10.1016/j.compositesb.2014.09.005. [6] R.R. Kieschke, T.W. Clyne, Development of a diffusion barrier for SiC monofilaments in titanium, Mater. Sci. Eng. A 135 (1991) 145–149, https://doi.org/10. 1016/0921-5093(91)90552-X. [7] Y.C. Fu, N.L. Shi, D.Z. Zhang, R. Yang, Effect of C coating on the interfacial microstructure and properties of SiC fiber-reinforced Ti matrix composites, Mater. Sci. Eng. A 426 (2006) 278–282, https://doi.org/10.1016/j.msea.2006.04.029. [8] X. Luo, Y. Yang, Q. Sun, Y. Yu, B. Huang, Y. Chen, Effect of Cu/Mo duplex coating on the interface and property of SiCf/Ti-6Al-4V composite, Mater. Sci. Eng. A 535 (2012) 6–11, https://doi.org/10.1016/j.msea.2011.12.029. [9] R. Vancheeswaran, D.G. Meyer, H.N.G. Wadley, Optimizing the consolidation of titanium matrix composites, Acta Mater. 45 (1997) 4001–4018, https://doi.org/10. 1016/S1359-6454(97)00104-3. [10] M. Wu, K. Zhang, H. Huang, M. Wang, H. Li, S. Zhang, M. Wen, Interfacial reactions in SiCf/C/Ti17 composites dominated by texture of carbon coatings, Carbon 124 (2017) 238–249, https://doi.org/10.1016/j.carbon.2017.08.065. [11] M.-J. Wang, H. Huang, H. Li, S.-M. Zhang, M. Wen, C. Song, F. Pan, Microstructure and interfacial strength of SiC fiber-reinforced Ti17 alloy composites with different consolidation temperatures, Rare Met. 37 (2018) 759–768, https://doi.org/10. 1007/s12598-018-1002-5. [12] S. Mukherjee, C.R. Ananth, N. Chandra, Effects of interface chemistry on the fracture properties of titanium matrix composites, Compos. Appl. Sci. Manuf. 29 (1998) 1213–1219, https://doi.org/10.1016/S1359-835X(97)00129-2. [13] X. Luo, Y. Yang, Y. Yu, X. Wang, B. Huang, Y. Chen, Effect of Mo coating on the interface and mechanical properties of SiC fiber reinforced Ti-6Al-4V composites, Mater. Sci. Eng. A 550 (2012) 286–292, https://doi.org/10.1016/j.msea.2012.04. 072. [14] E. Olevsky, H.J. Dudek, W.A. Kaysser, HIPing conditions for processing of metal matrix composites using the continuum theory for sintering—I. Theoretical analysis, Acta Mater. 44 (1996) 707–713, https://doi.org/10.1016/1359-6454(95) 00179-4. [15] E. Olevsky, H.J. Dudek, W.A. Kaysser, Hiping conditions for processing of metal matrix composites using continuum theory for sintering—II. Application to fibre reinforced titanium alloys, Acta Mater. 44 (1996) 715–724, https://doi.org/10. 1016/1359-6454(95)00180-8. [16] Z. Zhang, J. Qin, Z. Zhang, Y. Chen, W. Lu, D. Zhang, Effect of β heat treatment temperature on microstructure and mechanical properties of in situ titanium matrix composites, Mater. Des. 31 (2010) 4269–4273, https://doi.org/10.1016/j.matdes. 2010.04.003. [17] S. Schüler, B. Derby, M. Wood, C. Ward-Close, Matrix flow and densification during the consolidation of matrix coated fibres, Acta Mater. 48 (2000) 1247–1258, https://doi.org/10.1016/S1359-6454(99)00428-0.
[18] G.M. Zhao, Y.Q. Yang, W. Zhang, X. Luo, B. Huang, Y. Chen, Microstructure and grain growth of the matrix of SiCf/Ti–6Al–4V composites prepared by the consolidation of matrix-coated fibers in the β phase field, Compos. B Eng. 52 (2013) 155–163, https://doi.org/10.1016/j.compositesb.2013.04.024. [19] P.R. Subramanian, S. Krishnamurthy, S.T. Keller, M.G. Mendiratta, Processing of continuously reinforced Ti-alloy metal matrix composites (MMC) by magnetron sputtering, Mater. Sci. Eng. A 244 (1998) 1–10, https://doi.org/10.1016/S09215093(97)00820-4. [20] H.X. Peng, F.P.E. Dunne, K.H. Baik, P.S. Grant, Fibre re-arrangement and matrix softening phenomena in matrix-coated fibre (MCF) composites during vacuum hot pressing consolidation, Mater. Sci. Eng. A 346 (2003) 246–253, https://doi.org/10. 1016/S0921-5093(02)00540-3. [21] Q. Sun, X. Luo, Y.Q. Yang, G.H. Feng, G.M. Zhao, B. Huang, A review on the research progress of push-out method in testing interfacial properties of SiC fiberreinforced titanium matrix composites, Compos. Interfac. 22 (2015) 367–386, https://doi.org/10.1080/09276440.2015.1032150. [22] O. Knotek, F. Löffler, L. Wolkers, Phase stability and formation of Ti-Al-C-N PVD coatings, Surf. Coating. Technol. 68–69 (1994) 176–180, https://doi.org/10.1016/ 0257-8972(94)90156-2. [23] Y. Chen, M. Chu, L. Wang, X. Bao, Y. Lin, J. Shen, First‐principles study on the structural, phonon, and thermodynamic properties of the ternary carbides in Ti–Al–C system, Phys. Status Solidi 208 (2011) 1879–1884, https://doi.org/10. 1002/pssa.201127020. [24] S.Q. Guo, Y. Kagawa, H. Saito, C. Masuda, Microstructural characterization of interface in SiC fiber-reinforced Ti–15V–3Cr–3Al–3Sn matrix composite, Mater. Sci. Eng. A 246 (1998) 25–35, https://doi.org/10.1016/S0921-5093(97)00737-5. [25] D. Osborne, N. Chandra, H. Ghonem, Interphase behavior of titanium matrix composites at elevated temperature, Compos. Appl. Sci. Manuf. 32 (2001) 545–553, https://doi.org/10.1016/S1359-835X(00)00082-8. [26] A.M. Baker, P.S. Grant, M.L. Jenkins, The response of SiC fibres to vacuum plasma spraying and vacuum hot pressing during the fabrication of titanium matrix composites, J. Microsc. 196 (1999) 162–174, https://doi.org/10.1046/j.1365-2818. 1999.00625.x. [27] W. Zhang, Y.Q. Yang, G.M. Zhao, B. Huang, Z.Q. Feng, X. Luo, M.H. Li, J.H. Lou, Investigation of interfacial reaction in SiC fiber reinforced Ti–43Al–9V composites, Intermetallics 33 (2013) 54–59, https://doi.org/10.1016/j.intermet.2012.09.025. [28] Y.W. Xun, M.J. Tan, J.T. Zhou, Processing and interface stability of SiC fiber reinforced Ti–15V–3Cr matrix composites, J. Mater. Process. Technol. 102 (2000) 215–220, https://doi.org/10.1016/S0924-0136(00)00473-8. [29] K.L. Choy, Functionally graded coatings on SiC fibres for protection in Ti-based metal matrix composites, Scripta Mater. 34 (1996) 1753–1758, https://doi.org/10. 1016/1359-6462(96)00052-8. [30] S. Guo, C. Hu, H. Gao, Y. Tanaka, Y. Kagawa, SiC(SCS-6) fiber-reinforced Ti3AlC2 matrix composites: interfacial characterization and mechanical behavior, J. Eur. Ceram. Soc. 35 (2015) 1375–1384, https://doi.org/10.1016/j.jeurceramsoc.2014. 11.034. [31] H. Mei, Q. Bai, Y. Sun, H. Li, H. Wang, L. Cheng, The effect of heat treatment on the strength and toughness of carbon fiber/silicon carbide composites with different pyrolytic carbon interphase thicknesses, Carbon 57 (2013) 288–297, https://doi. org/10.1016/j.carbon.2013.01.076. [32] B. Huang, Y. Yang, H. Luo, M. Yuan, Y. Chen, Effect of the interfacial reaction layer thickness on the thermal residual stresses in SiCf/Ti–6Al–4V composites, Mater. Sci. Eng. A 489 (2008) 178–186, https://doi.org/10.1016/j.msea.2008.01.007.
17774
Contents lists available at ScienceDirect
Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Interfacial reactions and matrix microstructure evolution in SiCf/Ti composites dominated by primary structure of Ti matrix
T
Shuming Zhanga, Minjuan Wangb, Mao Wenc, Ming Wuc, Qingfeng Wanga,∗, Hao Huangb,∗∗ a
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, 066004, China AECC Beijing Institute of Aeronautical Materials, Beijing, 81-15, 100095, China c State Key Laboratory of Superhard Materials, Department of Materials Science, Key Laboratory of Automobile Materials, MOE, Jilin University, Changchun, 130012, China b
ARTICLE INFO
ABSTRACT
Keywords: SiCf/Ti composite Microstructure Interfacial reaction Strength
High-performance continuous SiC fiber reinforced titanium matrix composites (SiCf/Ti) are strongly required by aerospace vehicles, and their properties strongly depend on the microstructure of the interfacial reaction layer and the matrix. Thus, it is important to explore new routes to regulate these two SiCf/Ti components. In this work, α-phase-dominated and β-phase-dominated primary PVD Ti17 coatings were fabricated by controlling the bias voltage during physical vapor deposition (PVD). Their effects on the matrix microstructure evolution and interfacial reactions in SiCf/Ti after consolidation were further explored. It was found that compared with a primary α-Ti PVD Ti17 coating, a primary β-Ti PVD coating induced competitive growth between α-Ti and β-Ti grains in SiCf/Ti17 during consolidation. This yielded a smaller grain size and more β phase in the SiCf/Ti17 evolved from the primary β-Ti coating. In addition, a thicker reaction layer formed, as well as a lower interface strength and tensile strength in the SiCf/Ti17 evolved from a primary β-Ti coating, compared with the SiCf/Ti17 evolved from a primary α-Ti PVD Ti17 coating.
1. Introduction Light-weight high-strength structural materials, especially those possessing high specific strength and stiffness at both room and elevated temperatures, are highly desired for the construction of aerospace vehicles and advanced propulsion systems. Continuous SiC fiber reinforced titanium matrix composites (SiCf/Ti) have been explored as potential candidates in such applications [1–3]. It is generally accepted that interfaces play a crucial role in determining the properties of SiCf/ Ti, such as tensile strength and fatigue failure [4,5]. Thus, many researchers have investigated and tailored interfacial behaviors of SiCf/Ti to improve their performance via barrier layer optimization and adjustment of consolidation processes [6–9]. Wu et al. improved the tensile strength of SiCf/Ti through optimization of its interfacial properties by controlling the texture of the carbon coating, in which the distribution of interfacial reaction zones evolved from continuous elemental diffusion and grain growth during consolidation process was governed by the texture of the C coatings, consisting of turbostratic amorphous carbon, fine-grained carbon, transition TiC, coarse-grained TiC, and TiC with a decreasing C content gradient [10]. In addition, by
∗
adjusting the hot isostatic pressing (HIP)consolidation temperature, Wang et al. observed that the reduce or disappearance of fine TiCx layer thickness but rapid increment of the thickness of total reaction layers with elevating the consolidation temperature [11]. Barrier coatings and consolidation processes are commonly applied to control the thickness and microstructure of interfacial reactions, allowing the mechanical properties of SiCf/Ti composites to be further tailored. Since the interfacial microstructure strongly influences the properties of SiCf/Ti [12,13], there is a need to explore how interfacial reactions are affected by other avenues. Notably, during the deformation and failure of SiCf/Ti, the matrix usually provides nucleation sites for dislocations and cracks, and the subsequent expansion to interface zones can induce interfacial debonding [14,15]. This means that in addition to the interfacial microstructure, the matrix may also play an important role in influencing the mechanical properties of SiCf/Ti. Some recent attempts have been made to regulate the matrix microstructure by changing the consolidation temperature and duration [16–18]. The matrix-coated fiber (MCF) method has been widely utilized to prepare high-performance SiCf/Ti over the past several few years due to its unique advantages [19,20], in
Corresponding author. Corresponding author. E-mail addresses: [email protected] (Q. Wang), [email protected] (H. Huang).
∗∗
https://doi.org/10.1016/j.ceramint.2019.05.347 Received 13 March 2019; Received in revised form 27 May 2019; Accepted 30 May 2019 Available online 31 May 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
which a Ti alloy coating is deposited onto SiC fibers by physical vapor deposition (PVD) to serve as the primary structure of matrix. Obviously, PVD Ti alloy, directly contacted with SiC fibers, would become as both the reaction front and matrix source whose microstructure would strongly influence the interfacial reaction and matrix microstructure. However, the interfacial reaction and matrix microstructure in the SiCf/ Ti dominated by primary microstructure of PVD Ti alloy is scarcely investigated. In this paper, α-phase-dominated and β-phase-dominated primary PVD Ti17 coatings were prepared by controlling the bias voltage during PVD, and their effects on the interfacial reaction and matrix microstructure evolution in the SiCf/Ti after consolidation were also explored. The results showed that the primary structure of the PVD Ti17 coating determined both the interfacial reaction and matrix microstructure of SiCf/Ti. This work can provide valuable insight into the influence of the PVD Ti alloy structure on the matrix microstructure and interfacial reaction to further improve the performance of the SiCf/Ti. 2. Experimental procedures and characterization 2.1. Sample preparation The MCF method was used to fabricate Ti17 coatings onto tungstencored SiC monofilaments as precursor wires in a facing-targets magnetron-sputtering system, in which the Ti17 alloy (nominal composition in wt.%: 4.5–5.5 Al, 1.6–2.4 Sn, 1.6–2.4 Zr, 3.5–4.5 Mo, 3.5–4.5 Cr, and balance Ti) was utilized as the target, and C-coated SiC fibers with ∼100 μm diameters were used as the substrate. By regulating bias voltage, both α-phase- and β-phase-dominated PVD Ti coatings were successfully fabricated, and used as the primary matrix structure to be further evolved into an SiCf/Ti matrix after consolidating. Next, these precursor wires were placed into a Ti17 alloy canister hermetically to conduct sealing process using electron beam wielding. Subsequently, SiCf/Ti17 was manufactured by consolidating these packed precursor wires which have been sealed in a canister through HIP under the condition of 920 °C/120MPa/120mins, and then the furnace was cooled to room temperature. After HIP, both α-phase- and β-phase-dominated SiCf/Ti17 precursor wires, named as SiCf/α-Ti17 and SiCf/β-Ti17 respectively, were intact and compacted as shown in Fig. 1, and exhibited highly uniform distributions of SiC fibers. 2.2. Characterization Tensile samples were prepared for each HIP consolidated specimen with gauge dimensions of Φ3 × M6. Room tensile tests were performed at an extension rate of 1 mm/min, using an Inspekt Table 100 kN model universal testing machine. PVD titanium alloy was sputtered on a single-crystal silicon wafer, and one piece was peeled off for microstructure characterization. Primary Ti matrices were observed in parallel with the direction of columnar crystal growth. Both X-ray diffraction (XRD, D/max-2500) and field emission JEOL 2010F transmission electron microscopy (TEM) were employed to characterize
PVD Ti17 coatings to identify the primary phase structure of the matrix. The TEM specimens were thinned by a Leica EM RES 102 ion beam milling system. The interfacial reaction products were investigated by TEM, and the chemical compositions of each sub-layer in the interfacial zones were further examined by energy dispersive X-ray spectroscopy (EDS) equipped to the TEM. TEM specimens were cut from the composites perpendicular to the cross-section and in the radial direction of the fiber by FEI Quanta 200 FEG focused ion beam (FIB) milling in the process shown in Fig. 2. Scanning electron microscopy (SEM, FEI nano 450) was utilized to obtain micrographs of the samples. SiCf/Ti17 specimens used for SEM observations were cut along the cross-section, polished, and etched in an HF solution. The interfacial debonding strength was evaluated by performing fiber push-out tests on the sample slices with 0.5 ± 0.01 mm thicknesses [21]. Fractographies were observed by SEM and laser scanning confocal microscopy (LSCM, KEYENCE VK-100). 3. Results and discussion The nanocrystalline nature of PVD Ti alloys, containing large numbers of grain boundaries and different phase structures, would significantly affect the interfacial reactions and matrix microstructure of SiCf/Ti during consolidation. Two typical phase structures, including α-Ti nanocrystalline and β-Ti nanocrystalline, can be first obtained in the primary PVD Ti coating by regulating the PVD conditions to investigate its effects on the interfacial reactions and matrix evolution of SiCf/Ti. 3.1. Microstructure of primary PVD Ti17 coating As expected, α-Ti and β-Ti structures were obtained in the primary PVD Ti17 coatings fabricated at −150 V and −250 V of bias voltage, respectively, as shown by XRD and selected area electron diffraction (SAED)results. The Ti17 coatings were first analyzed by XRD, and the resulting spetra are shown in Fig. 3. At −150 V, a strong peak located at 36.2° appears, corresponding to hcp α-Ti 101¯0 , a stable structure at room temperature. Conversely, at −250 V, the α-Ti 101¯0 peak disappeared, and a weak and broad peak centered at 37.6° appeared, which was attributed scribed to bcc β-Ti 110 . This implies that besides the crystal structure, the high bias voltage of −250 V also reduced the crystallinity or created smaller grain sizes relative to −150 V due to the higher incident energy, which induced re-nucleation. The average grain size was estimated from the full-width half maximum in Fig. 3 to be ∼15 nm for −150 V and ∼4 nm for −250 V. Moreover, due to the higher compressive stress introduced by the bombardment of higherenergy particles (−250 V), the metastable β-Ti nucleation and growth was stabilized [22]. The SAED pattern in Fig. 4 further suggested that an α-Ti structure formed at −150 V and a β-Ti structure formed at −250 V, by assigning all patterns in −150 V and −250 V to α-Ti and βTi, respectively. The HRTEM images in Fig. 4 show typical polycrystalline character for both coatings, in which the grain size is 10–20 nm for −150 V and 3–8 nm for −250 V accompanied by partial amorphous zones, in agreement with the XRD results. The cross-sectional morphologies of the precursor wires were also investigated by SEM (Fig. 4b and d). Low-energy particle bombardment at −150 V resulted in typical column growth due to the shadow effect, whereas higher-energy energy particle bombardment at −250 V promoted the formation of dense structures due to higher diffusion and ion sub-implantation effects. 3.2. Microstructure of SiCf/Ti17
Fig. 1. Cross-sectional SEM micrographs of (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17.
3.2.1. Microstructure of titanium alloy matrix After being consolidated at the same HIP condition (in the α+β phase fields), the fabricated α-Ti and β-Ti Ti17 coatings on SiC fibers mainly evolved into corresponding SiCf/α-Ti17 and SiCf/β-Ti17
17768
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
Fig. 2. The process used to cut SiCf/α-Ti17 samples by FIB for (a)–(d).
Fig. 3. XRD patterns of primary PVD Ti17 coatings.
matrices in addition to forming some interfacial zones. The effect of the primary phase structure of PVD Ti on the interfacial reaction and matrix microstructure of SiCf/α-Ti17 and SiCf/β-Ti17 is clarified as follows. The representative matrix microstructures of SiCf/α-Ti17 and SiCf/βTi17 respectively evolved from primary α-Ti and β-Ti observed by SEM, as shown in Fig. 5. It was found that both composites exhibited typical equiaxed structures, but there were large differences in the grain sizes and content of each phase (α phase and β phase). The grain size of equiaxed α phase in the SiCf/α-Ti17 was much larger than in SiCf/βTi17, and SiCf/α-Ti17 also contained less β phase than SiCf/β-Ti17 (34% β phase for SiCf/α-Ti17 and 43% for SiCf/β-Ti17). The above differences in grain size and phase content in the SiCf/α-Ti17 and SiCf/ β-Ti17 mainly arose from the different growth behaviors during consolidation. Notably, as the α+β phase fields were applied to HIP
consolidation, both continuous grain growth and temperature-induced phase transition of matrix simultaneously occurred during the heating and isothermal stages. During the heating stage of SiCf/β-Ti17 consolidation, the primary metastable β-Ti phase became unstable and partially transformed into stable α-Ti, and the newly formed α-Ti coexisted with the existing β-Ti to impede the grain growth of each phase. In contrast, during the consolidation of SiCf/α-Ti17, the primary stable α-Ti phase tended to continually grow without phase transition before reaching the temperature of α+β phase field. That is to say, the competitive growth between α-Ti and β-Ti grains in SiCf/β-Ti17 during heating gave rise to smaller grains and more β phase before entering into the isothermal stage, compared with SiCf/α-Ti17. At the isothermal stage (α+β phase field), the α phase began to partially transform into β phase, and each phase continued to grow. Therefore, compared with SiCf/α-Ti17, the remarkable competitive grain growth between α-Ti and β-Ti grains in SiCf/β-Ti17 during heating resulted in much finer equiaxed α-Ti grains. Meanwhile, residual primary β-Ti grains in SiCf/ β-Ti17 also contributed to more β phase. 3.2.2. Microstructure of interfacial reaction layer Besides the matrix structure of SiCf/Ti, the primary phase structure of PVD Ti17 also strongly influenced the evolution of interfacial reaction layer (RL) during HIP consolidation. From the SEM images in Fig. 5, it is observed that RL in both SiCf/α-Ti17 and SiCf/β-Ti17, located between the C coating and Ti17 matrix, were compact without any cracks or holes, but a thicker RL was observed in SiCf/β-Ti17 (∼1.7 μm for SiCf/α-Ti17 and ∼2.4 μm for SiCf/β-Ti17). It is accepted that the RL mainly arose from the continuous diffusion of elemental carbon from the carbon coating into the adjacent Ti matrix and to form TiCx products during high-temperature consolidation, known as Ti + C = TiC [23,24]. The distribution of reaction products in the RL was confirmed to play an important role on the interfacial properties
17769
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
Fig. 4. The HRTEM images and SAED patterns of (a) SiCf/α-Ti17 and (c) SiCf/β-Ti17; the cross-sectional images of Ti17 precursor wires of (b) SiCf/α-Ti17 and (d) SiCf/β-Ti17.
Fig. 5. Microstructure of SiCf/Ti17 observed by SEM for (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17.
[12,25]. To clarify the product distribution and PVD Ti structuredominated interfacial reactions, the interfacial reaction zones were further investigated by TEM. The TEM images are shown in Fig. 6. As determined using SAED patterns, only the TiCx phase that was found in the RL of both composites was located between the C coating and the Ti17 matrix. However, the grain size of TiCx products showed a broad distribution, ranging from several nanometers near the C coating to hundreds of nanometers or more close to the Ti17 matrix, nearly approaching the grain size of the Ti17 matrix. Based on the TiC grain sizes, the RL can be divided into three sub-layers, namely Ⅰ RL, Ⅱ RL, and Ⅲ RL, which respectively correspond to fine-grained layers adjacent to the C coating, the transition layer in the middle of the RL zone, and the coarse-grained layer next to the matrix [10]. The corresponding thicknesses and grain sizes of the three sub-layers in the RL region and grain size of the Ti matrix for the two samples are listed in Table 1. For SiCf/α-Ti17, the total thickness of RL was determined to be∼1.7 μm,
consisting of ∼60 nm Ⅰ RL, ∼200 nm Ⅱ RL, and∼1.4 μm Ⅲ RL. Meanwhile, the grain size rapidly increased from less than 20 nm in Ⅰ RL, to ∼200 nm in Ⅱ RL, and larger than 500 nm in Ⅲ RL. The rapid increase in grain size along RL can be explained by noting that C atoms continuously diffused from the C coating towards the matrix, whose grains also grew during HIP. It has been previously confirmed that amorphous C usually exists in Ⅰ RL near the C coating due to the rich C source [26,27]. It impedes the TiC grain growth and maintains fine TiC grains (several nanometers), supported by the EDS result that C is abundant in Ⅰ RL (Table 2). It is clear that Ⅰ RL forms during the initial HIP stage, in which the C-rich character and nanocrystalline-boundaries can provide short-circuit paths for the diffusion of C atoms [28], which is described as a reaction-controlled process [1]. During the formation of Ⅰ RL, grain growth simultaneously occurred at zones near Ⅰ RL. As the reaction continues, the entry of C atoms across Ⅰ RL yields Ⅱ RL, which is responsible for the abrupt increase in grain size to ∼200 nm in Ⅱ RL. The newly formed Ⅱ RL with a much larger grain size exhibits nearstoichiometric character, as shown by the EDS result, and it can serve as a diffusion barrier layer [29]. Since the diffusion rate of C in the TiC lattice is a crucial route for C migration across Ⅱ RL, further interfacial reactions would be diffusion-controlled processes [1]. Obviously, the diffusivity rate of C in the TiC lattice is several orders of magnitude lower than in the Ti lattice [24] and is a short-circuit process. After C atoms diffuse across the Ⅰ RL and Ⅱ RL via further reactions, they would react with the grown Ti grains to form coarse-grained Ⅲ RL with grain sizes larger than 500 nm. The same variety of grain sizes along the RL also appeared in SiCf/βTi17, increasing from less than 20 nm in Ⅰ RL to larger than 200 nm in Ⅲ RL due to the combined roles of continuous C diffusion towards the matrix and the simultaneous grain growth of the matrix. The thickness of each sub-layer in SiCf/β-Ti17 was evaluated to be approximately 100 nm for Ⅰ RL, 400 nm for Ⅱ RL, and 1.9 μm for Ⅲ RL, which were ∼2.4 μm thicker than the RLs in SiCf/α-Ti17 (1.7 μm). Moreover, the grain size of each sublayer in SiCf/β-Ti17 was also much smaller than SiCf/α-Ti17, meaning that the primary α-Ti PVD precursor resulted in much larger grain sizes but thinner RLs compared with the primary β-Ti PVD precursor. Since the β-Ti PVD precursor was used as a reaction front, the additional phase transition from β-Ti to stable α-Ti occurred during interfacial reactions, yielding smaller matrix grains due to
17770
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
Table 2 EDS results of three sub-layers in the RL region of SiCf/α-Ti17 and SiCf/β-Ti17. SiCf/α-Ti17
Spot 1 (RL I) Spot 2 (RL II) Spot 3 (RL III)
SiCf/β-Ti17
C (at.%)
Ti (at.%)
C (at.%)
Ti (at.%)
64.27 42.13 28.05
28.36 55.98 67.15
59.77 38.36 28.93
34.86 61.22 69.67
RL- Reaction Layer.
Fig. 6. TEM brightfield images and SAED patterns of interfacial reaction products and Ti17 matrix for (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17.
competitive grain growth. After C diffused into the nearby matrix to form a RL, the smaller grain sizes also appeared in the RL of SiCf/βTi17. In addition, more β-Ti phase was present in the SiCf/β-Ti17 matrix relative to SiCf/α-Ti17, as revealed by the SEM images in Fig. 5. It has been previously reported that β titanium alloy matrices react more rapidly with SiC fiber than α titanium alloys due to the faster diffusion rate of C or Si in a β-Ti matrix [24]. Thus, the faster C diffusion in β-Ti
than α-Ti would be mainly responsible for the formation of a thicker RL. The smaller grain size also partially contributed to a faster diffusion rate due to more terrible crystalline-boundary diffusion. In the two samples, the elemental distribution of carbon and titanium along the growth direction of the reaction layer from spot 1 to spot 3, respectively, are given in Table 2. The data shows that Ⅰ RL is carbon rich, near-stoichiometric in Ⅱ RL, and C-poor in Ⅲ RL. The carbon content of each sample decreased as the distance away from C coating increased, and conversely increased with titanium, further supporting the above viewpoint of interdiffusion processes. The STEM images and EDS elemental maps were obtained to further visualize the elemental distribution of the interfacial zones of both SiCf/ α-Ti17 and SiCf/β-Ti17, as shown in Fig. 7. The matrix elements were almost completely blocked outside the C coating, but C atoms obviously diffused from the C coating towards the matrix and dominated the formation of interfacial reaction zones with gradient C content. Additionally, the clear boundaries of carbon coating, interfacial reaction zones, and titanium alloy matrix were distinguished in both SiCf/α-Ti17 and SiCf/β-Ti17, suggesting the same RL thickness obtained from the SEM images in Fig. 5. During HIP, the formation of interfacial reaction zones is mainly controlled by the continuous diffusion of C atoms from the C coating into the adjacent Ti matrix. The content and grain size of each phase in the Ti matrix near the C diffusion front also determines the C diffusion rate. The microstructure evolution of the matrix in primary α-Ti or β-Ti nanocrystalines was mainly dominated by the continuous grain growth and temperature-induced phase transition, as shown in Fig. 8. The main difference in the SiCf/α-Ti17 and SiCf/β-Ti17 matrices were due to the heating stage T1-T2, in which competitive growth between α-Ti and βTi grains occurred in SiCf/β-Ti17, yielding smaller grain sizes and more β phase than SiCf/α-Ti17 before entering T3. Consequently, the smaller grain size and additional β phase in SiCf/β-Ti17 provided more terrible crystalline-boundary diffusion and faster C diffusion in β-Ti relative to SiCf/α-Ti17. This resulted in the formation of a thicker RL and fewer and finer equiaxed α-Ti grains in SiCf/β-Ti17, compared with SiCf/αTi17, as illustrated in Fig. 8. 3.3. Interfacial strength and tensile strength Since primary α-Ti and β-Ti PVD coatings were used to fabricate SiCf/α-Ti17 and SiCf/β-Ti17, different interfacial RLs, including different grain sizes and thicknesses of each sublayer were observed in each of the composites. This would further enhance the interface strength of SiCf/α-Ti17 and SiCf/β-Ti17. It has been widely accepted that the interface strength of SiCf/Ti is a crucial parameter governing
Table 1 The corresponding thicknesses and grain sizes of three sub-layers in the reaction layer region and grain sizes of the Ti matrix for two samples. Thickness of RL
SiCf/α-Ti17 SiCf/β-Ti17
Grain size of TiC
Grain size of Ti matrix/μm
RL I/nm
RL II/nm
RL III/μm
RL I/nm
RL II/nm
RL III/nm
60 100
200 400
1.4 1.9
≤20 ≤20
= Thickness of RL II 20 ≤ TiC≤ 100
≥500 ≥200
RL- Reaction Layer. 17771
0.4–2 0.1–0.4
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
Fig. 7. STEM images and EDS elemental maps of the interfacial zones of (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17.
the failure modes of composites under external loads, and therefore has a significant influence on mechanical properties [30]. Thus, push-out experiments were carried out to evaluate the interfacial properties of SiCf/α-Ti17 and SiCf/β-Ti17, as shown in Fig. 9. In the load-displacement curves, the peak load Pd is considered to be the point at which the load corresponds to the onset of fiber protrusion from the back surface, and the average Pd of SiCf/α-Ti17 and SiCf/β-Ti17 was measured to be ∼17.5 N and ∼13.1 N, respectively. By substituting Pd values, the shear stress or interface strength τd can be calculated based on the following formula [31]: d
= Pd/ df h
(1)
where df is the diameter of the SiC fiber and h is the thickness of the composite. The calculated τd was ∼111 MPa for SiCf/α-Ti17 and ∼82 MPa for SiCf/β-Ti17, which approach the reported values in other systems measured using the same method. Fu et al. reported a τd of ∼118 MPa for an SiCf/C/Ti–6Al–4V system [7] and D. Osborne et al.
Fig. 9. Typical fiber push-out load-displacement curves for SiCf/α-Ti17 and SiCf/β-Ti17.
Fig. 8. Models of titanium alloy crystal growth of different structure PVD precursors during heat treatment, showing how it affects the C diffusion and reaction interface formation.
17772
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al.
in the RL of SiCf/β-Ti17. It is clear that the SiCf/α-Ti17 with a thinner RL exhibited a higher interfacial strength than SiCf/β-Ti17, and when RL thickness of SiCf/β-Ti17 reached 2.4 μm it may have promoted preferential crack nucleation at the RL Tensile tests were performed on SiCf/α-Ti17 and SiCf/β-Ti17, and the resulting curves are shown in Fig. 11. The average tensile strengths of SiCf/α-Ti17 and SiCf/β-Ti17 were calculated to be 2094 MPa and 1910 MPa, respectively, and each composite had a much higher tensile strength than the Ti17 matrix (1060 MPa). In the composite, both the SiC fibers and the interfacial properties significantly affected the tensile strength. The tensile strength of SiCf/α-Ti17 was ∼200 MPa higher than SiCf/β-Ti17 despite having the same fiber volume fractions. To further analyze this result, the fracture morphologies of SiCf/α-Ti17 and SiCf/β-Ti17 were observed as shown in Fig. 12. Both composites show typical step-like fractures involving obvious push-out phenomena of their short fibers, rather than either brush-like fracture caused by a too-low τd or brittle facture induced by a too-high τd due to lower interfacial debonding. The morphologies suggested that each composite should possess suitable interface strength, as evidenced by their high tensile strengths (> 1900 MPa). More detailed observations showed that the lower τd of SiCf/β-Ti17 resulted in the appearance of obviously longer push-out fibers than SiCf/α-Ti17, which agreed with classical shear-lag model predictions. This may be responsible for the lower tensile strength of SiCf/β-Ti17 since it has a lower τd than SiCf/α-Ti17. Therefore, by only controlling the primary structure of PVD Ti17 coatings (α-Ti or β-Ti), the interfacial reactions and microstructure of SiCf/Ti17 can be tailored to achieve a higher τd and tensile strength in SiCf/α-Ti17 than SiCf/β-Ti17.
Fig. 10. SEM micrographs of interfacial cracks in SiCf/β-Ti17 for (a) and (b).
Fig. 11. Typical tensile stress–strain curves of SiCf/α-Ti17 and SiCf/β-Ti17.
4. Conclusions Both α-phase-dominated and β-phase-dominated primary PVD Ti17 coatings were successfully fabricated onto SiC fibers by only regulating the bias voltage, followed by consolidation into SiCf/α-Ti17 and SiCf/βTi17 via HIP. Based on the microstructure observations and mechanical properties of the two composites, the following conclusions can be drawn: 1) During HIP, competitive growth between α-Ti and β-Ti grains appeared in the SiCf/β-Ti17 during the heating stage because primary β-Ti grains became unstable, yielding smaller grain sizes and more β phase in SiCf/β-Ti17 than SiCf/α-Ti17. Compared with SiCf/β-Ti17, the smaller grain size and higher β phase in the matrix of SiCf/βTi17 also induced more terrible crystalline-boundary diffusion and faster C diffusion in β-Ti. This resulted in the formation of thicker RLs consisting of a fine-grained layer, a transition layer, and a coarse-grained layer. 2) SiCf/α-Ti17 with a thinner RL exhibited a higher interface strength than SiCf/β-Ti17, when the RL thickness of SiCf/β-Ti17 reached 2.4 μm which may have promoted preferential crack nucleation at the RL. The lower interfacial strength of SiCf/β-Ti17 caused increased fiber push-out than in SiCf/α-Ti17, contributed to the lower tensile strength in SiCf/β-Ti17. This provides a new strategy to regulate the matrix microstructure and RL of SiCf/Ti to further tailor their properties.
Fig. 12. Fracture morphologies of (a) SiCf/α-Ti17 and (b) SiCf/β-Ti17 as observed by SEM; fracture morphologies of (c) SiCf/α-Ti17 and (d) SiCf/β-Ti17 observed by LSCM.
observed a slightly higher τd of ∼165 MPa in an SCS-6/Timetal-21S system [25]. Even if same test method was employed to characterize τd many factors such as thickness, polishing, and the material systems of the tested samples would affect these values. Thus, it is difficult to quantitatively determine the optimum value of τd for SiCf/Ti composite systems, but a moderate τd is strongly preferred to effectively transfer the load of the matrix to SiC fibers and also to facilitate interfacial debonding instead of across the fibers [30]. When τd is too low, a brushlike fracture mode occurs, while a too high τd induces typical flat brittle mode fracture [13]. It is generally believed that the τd of composites can be determined by the RL thickness. RL with appropriate thicknesses can improve the τd, but above a critical thickness, τd would be worsened due to preferential crack nucleation occurring at the brittle RL during loading [32]. Fig. 10 shows the appearance of interfacial cracks
Acknowledgements This work was supported by the National Key Research and Development Program of China (Grant No. 2016YFB0301203) and the National Natural Science Foundation of China (Grant No. 51871109). References
17773
[1] Z.X. Guo, B. Derby, Solid-state fabrication and interfaces of fibre reinforced metal matrix composites, Prog. Mater. Sci. 39 (1995) 411–495, https://doi.org/10.1016/
Ceramics International 45 (2019) 17767–17774
S. Zhang, et al. 0079-6425(95)00002-X. [2] J.M. Larsen, S.M. Russ, J.W. Jones, An evaluation of fiber-reinforced titanium matrix composites for advanced high-temperature aerospace applications, Metall. Mater. Trans. A 26 (1995) 3211–3223, https://doi.org/10.1007/BF02669450. [3] C. Leyens, F. Kocian, J. Hausmann, W.A. Kaysser, Materials and design concepts for high performance compressor components, Aero. Sci. Technol. 7 (2003) 201–210, https://doi.org/10.1016/S1270-9638(02)00013-5. [4] Y.Q. Yang, H.J. Dudek, J. Kumpfert, Interfacial reaction and stability of SCS-6 SiC/ Ti–25Al–10Nb–3V–1Mo composites, Mater. Sci. Eng. A 246 (1998) 213–220, https://doi.org/10.1016/S0921-5093(97)00749-1. [5] G.H. Feng, Y.Q. Yang, X. Luo, J. Li, B. Huang, Y. Chen, Fatigue properties and fracture analysis of a SiC fiber-reinforced titanium matrix composite, Compos. B Eng. 68 (2015) 336–342, https://doi.org/10.1016/j.compositesb.2014.09.005. [6] R.R. Kieschke, T.W. Clyne, Development of a diffusion barrier for SiC monofilaments in titanium, Mater. Sci. Eng. A 135 (1991) 145–149, https://doi.org/10. 1016/0921-5093(91)90552-X. [7] Y.C. Fu, N.L. Shi, D.Z. Zhang, R. Yang, Effect of C coating on the interfacial microstructure and properties of SiC fiber-reinforced Ti matrix composites, Mater. Sci. Eng. A 426 (2006) 278–282, https://doi.org/10.1016/j.msea.2006.04.029. [8] X. Luo, Y. Yang, Q. Sun, Y. Yu, B. Huang, Y. Chen, Effect of Cu/Mo duplex coating on the interface and property of SiCf/Ti-6Al-4V composite, Mater. Sci. Eng. A 535 (2012) 6–11, https://doi.org/10.1016/j.msea.2011.12.029. [9] R. Vancheeswaran, D.G. Meyer, H.N.G. Wadley, Optimizing the consolidation of titanium matrix composites, Acta Mater. 45 (1997) 4001–4018, https://doi.org/10. 1016/S1359-6454(97)00104-3. [10] M. Wu, K. Zhang, H. Huang, M. Wang, H. Li, S. Zhang, M. Wen, Interfacial reactions in SiCf/C/Ti17 composites dominated by texture of carbon coatings, Carbon 124 (2017) 238–249, https://doi.org/10.1016/j.carbon.2017.08.065. [11] M.-J. Wang, H. Huang, H. Li, S.-M. Zhang, M. Wen, C. Song, F. Pan, Microstructure and interfacial strength of SiC fiber-reinforced Ti17 alloy composites with different consolidation temperatures, Rare Met. 37 (2018) 759–768, https://doi.org/10. 1007/s12598-018-1002-5. [12] S. Mukherjee, C.R. Ananth, N. Chandra, Effects of interface chemistry on the fracture properties of titanium matrix composites, Compos. Appl. Sci. Manuf. 29 (1998) 1213–1219, https://doi.org/10.1016/S1359-835X(97)00129-2. [13] X. Luo, Y. Yang, Y. Yu, X. Wang, B. Huang, Y. Chen, Effect of Mo coating on the interface and mechanical properties of SiC fiber reinforced Ti-6Al-4V composites, Mater. Sci. Eng. A 550 (2012) 286–292, https://doi.org/10.1016/j.msea.2012.04. 072. [14] E. Olevsky, H.J. Dudek, W.A. Kaysser, HIPing conditions for processing of metal matrix composites using the continuum theory for sintering—I. Theoretical analysis, Acta Mater. 44 (1996) 707–713, https://doi.org/10.1016/1359-6454(95) 00179-4. [15] E. Olevsky, H.J. Dudek, W.A. Kaysser, Hiping conditions for processing of metal matrix composites using continuum theory for sintering—II. Application to fibre reinforced titanium alloys, Acta Mater. 44 (1996) 715–724, https://doi.org/10. 1016/1359-6454(95)00180-8. [16] Z. Zhang, J. Qin, Z. Zhang, Y. Chen, W. Lu, D. Zhang, Effect of β heat treatment temperature on microstructure and mechanical properties of in situ titanium matrix composites, Mater. Des. 31 (2010) 4269–4273, https://doi.org/10.1016/j.matdes. 2010.04.003. [17] S. Schüler, B. Derby, M. Wood, C. Ward-Close, Matrix flow and densification during the consolidation of matrix coated fibres, Acta Mater. 48 (2000) 1247–1258, https://doi.org/10.1016/S1359-6454(99)00428-0.
[18] G.M. Zhao, Y.Q. Yang, W. Zhang, X. Luo, B. Huang, Y. Chen, Microstructure and grain growth of the matrix of SiCf/Ti–6Al–4V composites prepared by the consolidation of matrix-coated fibers in the β phase field, Compos. B Eng. 52 (2013) 155–163, https://doi.org/10.1016/j.compositesb.2013.04.024. [19] P.R. Subramanian, S. Krishnamurthy, S.T. Keller, M.G. Mendiratta, Processing of continuously reinforced Ti-alloy metal matrix composites (MMC) by magnetron sputtering, Mater. Sci. Eng. A 244 (1998) 1–10, https://doi.org/10.1016/S09215093(97)00820-4. [20] H.X. Peng, F.P.E. Dunne, K.H. Baik, P.S. Grant, Fibre re-arrangement and matrix softening phenomena in matrix-coated fibre (MCF) composites during vacuum hot pressing consolidation, Mater. Sci. Eng. A 346 (2003) 246–253, https://doi.org/10. 1016/S0921-5093(02)00540-3. [21] Q. Sun, X. Luo, Y.Q. Yang, G.H. Feng, G.M. Zhao, B. Huang, A review on the research progress of push-out method in testing interfacial properties of SiC fiberreinforced titanium matrix composites, Compos. Interfac. 22 (2015) 367–386, https://doi.org/10.1080/09276440.2015.1032150. [22] O. Knotek, F. Löffler, L. Wolkers, Phase stability and formation of Ti-Al-C-N PVD coatings, Surf. Coating. Technol. 68–69 (1994) 176–180, https://doi.org/10.1016/ 0257-8972(94)90156-2. [23] Y. Chen, M. Chu, L. Wang, X. Bao, Y. Lin, J. Shen, First‐principles study on the structural, phonon, and thermodynamic properties of the ternary carbides in Ti–Al–C system, Phys. Status Solidi 208 (2011) 1879–1884, https://doi.org/10. 1002/pssa.201127020. [24] S.Q. Guo, Y. Kagawa, H. Saito, C. Masuda, Microstructural characterization of interface in SiC fiber-reinforced Ti–15V–3Cr–3Al–3Sn matrix composite, Mater. Sci. Eng. A 246 (1998) 25–35, https://doi.org/10.1016/S0921-5093(97)00737-5. [25] D. Osborne, N. Chandra, H. Ghonem, Interphase behavior of titanium matrix composites at elevated temperature, Compos. Appl. Sci. Manuf. 32 (2001) 545–553, https://doi.org/10.1016/S1359-835X(00)00082-8. [26] A.M. Baker, P.S. Grant, M.L. Jenkins, The response of SiC fibres to vacuum plasma spraying and vacuum hot pressing during the fabrication of titanium matrix composites, J. Microsc. 196 (1999) 162–174, https://doi.org/10.1046/j.1365-2818. 1999.00625.x. [27] W. Zhang, Y.Q. Yang, G.M. Zhao, B. Huang, Z.Q. Feng, X. Luo, M.H. Li, J.H. Lou, Investigation of interfacial reaction in SiC fiber reinforced Ti–43Al–9V composites, Intermetallics 33 (2013) 54–59, https://doi.org/10.1016/j.intermet.2012.09.025. [28] Y.W. Xun, M.J. Tan, J.T. Zhou, Processing and interface stability of SiC fiber reinforced Ti–15V–3Cr matrix composites, J. Mater. Process. Technol. 102 (2000) 215–220, https://doi.org/10.1016/S0924-0136(00)00473-8. [29] K.L. Choy, Functionally graded coatings on SiC fibres for protection in Ti-based metal matrix composites, Scripta Mater. 34 (1996) 1753–1758, https://doi.org/10. 1016/1359-6462(96)00052-8. [30] S. Guo, C. Hu, H. Gao, Y. Tanaka, Y. Kagawa, SiC(SCS-6) fiber-reinforced Ti3AlC2 matrix composites: interfacial characterization and mechanical behavior, J. Eur. Ceram. Soc. 35 (2015) 1375–1384, https://doi.org/10.1016/j.jeurceramsoc.2014. 11.034. [31] H. Mei, Q. Bai, Y. Sun, H. Li, H. Wang, L. Cheng, The effect of heat treatment on the strength and toughness of carbon fiber/silicon carbide composites with different pyrolytic carbon interphase thicknesses, Carbon 57 (2013) 288–297, https://doi. org/10.1016/j.carbon.2013.01.076. [32] B. Huang, Y. Yang, H. Luo, M. Yuan, Y. Chen, Effect of the interfacial reaction layer thickness on the thermal residual stresses in SiCf/Ti–6Al–4V composites, Mater. Sci. Eng. A 489 (2008) 178–186, https://doi.org/10.1016/j.msea.2008.01.007.
17774