Text/Pedestrian who wants to break his soul
As a new material, high-strength Qin alloy has the advantages of high specific strength, easy heat treatment and good corrosion resistance, and is used as a key structural material in aerospace, and has made outstanding contributions to the development of satellite launches, launch vehicles and missiles.
By controlling the content of β stable elements, metastable β alloys can be obtained, which have been more widely used in structural components of aerospace vehicles. In addition, only a trace amount of B and C need to be added to the matrix of the alloy to synthesize TiB whiskers and TiC particle reinforced phases in situ, which can enhance dislocation accumulation, help twin nucleation, and effectively refine the grains, reduce the size of the secondary α slats, etc., and improve the strength of the composite.
The equivalent of Ti-B20 alloy (Ti-3.5A1-5Mo-3V-2Cr-1Fe-2Zr-2Sn) is at the lowest range of metastable β alloy, which has a high aging strengthening effect and good strength-plasticity matching effect. Therefore, we selected the new composite material of TiB and TiC hybrid reinforced T-B20 for research, which provides further theoretical support for the application of the composite material.
Heat treatment is an important means to strengthen β alloy, and excellent combination of mechanical properties can be obtained with the help of reasonable heat treatment process. The in-situ reinforced phase in the matrix and the precipitated phase during heat treatment will have a significant impact on the properties of the alloy.
1. Solid materials and methods
Our experimental material is (TiB+TiC/T-B20 composite, and TiB and TiC with a volume fraction of 2.2% are added to the T-3.5A1-5M-3V-2Cr-1Fe-2Zr-2Sn (Ti-B20) matrix, and the specific chemical composition of the material is shown in Table 1.
Table 1 (TiB+TiC)/Ti-B20 composite chemical composition
The ingot of the composite material is obtained by vacuum self-consumption melting process, and then the Qin alloy forging rod is obtained by the five-fire forging process. The phase transition temperature of the alloy measured by metallographic method is 840+5 °C. The microstructure and XRD diagram of the original wrought alloy are shown in Figure 1.
Figure 1 Microstructure and XRD diagram of the original wrought alloy
It can be seen from Figure 1 that the forged alloy is mainly composed of a single β phase, accompanied by diffraction peaks of lower strength such as (100) and (102) of the phase, corresponding to the precipitation of a small number of α primary phases. Some long slabs and strips of primary phase appear fragmentation, and at the same time, many small α phases appear inside the grain.
The prepared wrought alloy was subjected to the descending two-stage aging heat treatment (HLDA, high-lowduplexaging) shown in Table 2 by box-type resistance furnace, and the high-temperature oxidation resistance coating solution was applied to the surface of the alloy before heat treatment to prevent oxidation on the surface.
Table 2 Heat treatment protocols
Characterization of the microstructure of wrought alloys and heat-treated alloys. Phase analysis was performed using the Empyrean X-ray diffractometer and microstructure was observed using the Quanta200FEG scanning electron microscope (SEM). SEM specimens need to be prepared by water sandpaper polishing, electrolytic polishing and corrosion of corrosive fluid, and the ratio of corrosive liquid is 90vol.% distilled water, 7vol.% nitric acid, 3vo.% hydrofluoric acid.
The electronic universal testing machine (Instron-5500R) and RDL100 electron variable testing machine were used to test the tensile properties of the specimen after heat treatment at room temperature and the high temperature creep test of the alloy. The dimensions of the tensile specimen and creep specimen at room temperature are shown in Figure 2.
Fig. 2 Dimensions of tensile specimen (left) and creep specimen (right) at room temperature (unit: mm)
The stretching rate at room temperature is mm/min. The creep test conditions were 500°C/300MPa, 550C/250MPa, 550C/275MPa, 550C/300MPa, respectively. The temperature in the resistance furnace is heated to the set temperature and kept warm for a period of time before applying stress.
Second, the experimental results and discussion
2.1 Microstructure and mechanical properties of silver alloy after heat treatment
The wrought alloy was solved in the single-phase zone and the duplex zone, and then the two-stage aging treatment was carried out in descending order. The microstructure and XRD analysis results of the alloy after heat treatment are shown in Fig. 3 and Fig. 4.
Figure 3 Microstructure of an alloy after heat treatment
Figure 4 XRD diagram of an alloy after heat treatment
After the forged alloy was solved at 790C, long strips and convex lens-like primary phases were observed at the grain boundary and inside the grain, respectively. After the solid solution + descending two-stage aging treatment at 790°C, the intensity of the diffraction peaks of the α phase increased significantly, and new diffraction peaks such as (002), (103), (200), and (112) appeared, indicating that a significant β→α phase transition occurred in the descending aging, corresponding to the precipitation of a large number of small dispersed secondary α phases.
At the same time, there are many primary α phases of different shapes retained inside the alloy, which are spherical and convex lens-shaped, and randomly distributed in the matrix. The original continuous grain boundary α phase has undergone broken spheroidization, which is because in the phase transition process of β→α, Mo, V and other β stable elements are gathered in the β phase, AI and other α stable elements are concentrated in the α phase, and the concentration gradient generated between the two phases makes the elements diffuse each other at the phase interface, resulting in the continuous dissolution of the α phase, and the β phase is gradually wedged into the α phase, so that the grain boundary phase cracks and gradually spheroidizes under the action of surface tension.
Figure 5 Tensile stress-strain curves of wrought alloys and heat-treated alloys
Combined with the stress-strain curve of Figure 5 and the statistical results of mechanical properties in Table 3, it can be seen that compared with the wrought alloy, the tensile strength of the alloy after the two-stage aging treatment at 790 °C solid solution + descending order is slightly improved, from 1529MPa to 1565MPa, and the elongation is slightly reduced from 2.4% to 1.7%.
Table 3 Room temperature tensile properties of wrought alloys and alloys after heat treatment
On the one hand, the length-to-diameter ratio decreases after the spheroidization of the continuous grain boundary α phase, which shortens the effective distance of crack propagation along the grain boundary. On the other hand, the study of tensile deformation of Ti-5553 alloy by OINDY et al. shows that the primary α phase has sufficient plastic deformation ability, which has the effect of coordinating the deformation between sheet elements, which can improve the ductility of the alloy. In addition, the presence of these α primary phases will weaken the effect of secondary α phase dispersion enhancement, resulting in a small increase in strength and elongation.
After solid solution at 860 °C, the α phase dissolved and disappeared, and the structure was composed of a single β phase, and the TiB and TiC enhanced phases were clearly visible in the matrix. After the two-stage aging treatment of solid solution + descending order at 860°C, the inside of the grain is completely composed of fine secondary α phases, and the single microstructure characteristics greatly increase the tensile strength of the alloy to 1734MPa, while the elongation decreases to 0.8%. Numerous studies have shown that the larger the volume fraction of the secondary α phase, the greater the increase in strength and the decrease in plasticity of the alloy, which is consistent with our findings.
Fig. 6 Tensile fractures of wrought alloys and heat-treated alloys
Figure 6 shows the tensile fracture image of the original wrought alloy and the descending two-stage aging alloy.
The fracture surface of the original alloy is composed of a large number of fine and shallow tendons, which exhibit typical ductile fracture characteristics; TiC particle cracking and TiB whisker debonding phenomena are observed in some parts, and the reinforced phase is the main loading object during the stretching process, and stress concentration is easy to occur, resulting in cracking of the reinforcement and bonding separation of the reinforce/matrix interface, resulting in low plasticity of the alloy.
After the solid solution + descending two-stage aging at 790°C, the failure mode of the alloy remained unchanged, and it was still manifested as the ductile fracture mode of the concentrated load of the reinforced body. After 860°C solid solution + descending double-stage aging, the number and size of tendons decreased, some smooth fracture planes appeared, and the characteristics of brittle fractures increased, corresponding to their lower plasticity.
Variable properties of 2.2 state alloy after heat treatment
The high-temperature creep experiment was carried out on the alloy after the descending two-stage aging treatment, and the creep curve was first performed under the deformation condition of 500°C/300MPa, and the creep curve obtained was shown in Figure 7.
When the creep time reaches 60h, the experiment is stopped, and the creep has entered the steady-state creep stage. Compared with the alloy after solid solution + descending two-stage aging at 860 °C, the alloy with solid solution + descending two-stage aging at 790°C has a lower steady-state creep rate (2.54X10s and a smaller strain variable (8.86%).
Figure 7 Creep curve of alloy at 500°C/300MPa after heat treatment
Increase the creep temperature to 550 °C, and the creep curve and performance data of the alloy under different stress conditions are shown in Figure 8 and Table 4.
At this deformation temperature, the creep curves of the alloy after heat treatment show typical primary creep stage, steady-state creep stage and accelerated creep stage.
At the same temperature, with the increase of applied stress, the steady-state creep rate of the alloy after solid solution + descending two-stage aging at 790°C increased by 187% from 3.49x10s to 1.00x10, and the creep life decreased from 15.4h to 5.9h; the steady-state creep rate of the alloy after solid solution + descending two-stage aging at 860°C increased from 5.31x10s' to 1.25x10s, an increase of 135%, and the creep life decreased from 11.6h to 4.9h.
Under the same creep conditions, the alloy after 790°C solid solution + descending two-stage aging treatment has a lower steady-state creep rate and higher creep life, indicating that the alloy after 790°C solid solution + descending two-stage aging has better creep performance.
Figure 8 Creep curves of alloys in the range of 550C and 250-300 MPa after heat treatment
Table 4 Creep properties of alloys after heat treatment
The creep process of an alloy is controlled by two typical mechanisms: work hardening due to long-range dislocation action and strain recovery due to short-range dislocation action.
When work hardening and strain recovery reach dynamic equilibrium, it enters the creep steady-state stage, where the creep rate is the smallest, that is, the steady-state creep rate. The relationship between steady-state creep rate and temperature and stress is usually expressed by the Norton-Bailey rule:
Norton-Bailey rule
where ε̇m is the steady-state creep rate (s-1), A is the material constant, σ is the rheological stress, R is the ideal gas constant, and T is the absolute temperature (K),
n is the stress index and Qc is the creep apparent activation energy (kJ/mol).
According to the creep data in Table 4, the relationship between the steady-state creep rate and stress of the alloy after heat treatment can be obtained. Figure 9 shows the lnε̇-lnσ relationship diagram, according to the linear fitting calculation results, the stress indices of the alloy after solid solution + descending two-stage aging at 790°C/860C solid solution + descending order are 4.66 and 5.80 respectively.
According to related studies, at 4<n<7, the creep rate is mainly controlled by dislocation climbing19. Therefore, for alloys treated with solid solution + descending two-stage aging at different temperatures, the creep deformation mechanism is dislocation climbing.
Figure 9 Steady-state creep rate vs. stress plot of alloys
2.3 Creep microstructure and fracture characteristics of wrought alloy after heat treatment
The creep properties of the alloy after heat treatment are different, which is closely related to the heat treatment structure and creep parameters before creep. When the creep temperature is 550 °C, the creep microstructure under different stress conditions is shown in Figures 10 and 11.
Figure 10 Creep microstructure of alloy after HT1 treatment (550°C)
The pre-aging process of pre-aging materials precipitates secondary phases of different orientations, which can effectively hinder the dislocation movement and improve creep resistance. By comparing the microstructure before the change (Figure 3), it can be found that the morphology of the secondary phase distributed in the tissue after high temperature creep changes, and with the auxiliary action of applied stress, the phase undergoes obvious degradation and coarsening phenomenon, from the original fine needle-like feature to short and long feature and polygonal structure, which can be observed under different stresses.
Since the reduction of the length of the needle phase will increase the steady-state creep rate and reduce the creep resistance, the degradation phenomenon is detrimental to the creep performance. For the alloy treated by HTI, a large number of dislocations will occur inside the α phase during the creep deformation process, and the thermal diffusion ability of the dislocation is strong at a high temperature of 550C, so that a large number of primary α phases are dissolved.
A small number of micropores and cracks appear in the creep tissue, and the polymerization phenomenon of microcracks in the middle of adjacent TiB whiskers can be observed when the stress is 300MPa.
With the increase of applied stress, the thermal diffusion time is shortened, and the primary α phase retained in the tissue, especially the isostationary phase, gradually increases. However, the equiaxed α phase and the β matrix have a non-common interface, which has a high diffusion rate, which is not conducive to the improvement of variable resistance, which is consistent with the change trend of creep performance data in Table 4.
When the solution temperature is increased to 860 °C, the creep microstructure of the alloy becomes single, as shown in Figure 11, and fine dispersed secondary α phase structure is obtained under different stress conditions. A large number of micropores and microcracks are clearly visible, concentrated in the inside, both ends and adjacent areas of TiC particles and TiB whiskers, and the number of defects increases with the increase of stress.
Figure 11 Creep microstructure of the alloy after HT2 treatment (550°C)
The generation of micropores and cracks comes from the stress process and strengthening mechanism of TiC and TiB.
Firstly, the thermal expansion coefficients of pure Ti, TiB and TiC are 8.2×10-6K-1, 8.6×10-6K-1, (6.52~7.12)×10-6K-1, respectively, and there is a difference in the thermal expansion coefficient of TiC and the matrix, which will cause thermal mismatch strain during the stress, which will cause an increase in the dislocation density in the matrix near TiC, and play a certain dislocation strengthening effect.
The thermal expansion coefficient of TiB whisker and matrix is similar, and it has good combination with the matrix, and when its length-diameter ratio is greater than the critical length-diameter ratio, it can play a role in load transfer strengthening, and bear the load from the matrix through the interface. In addition, the C atom can also play a certain solid solution strengthening effect, strengthen the nailing effect on dislocations, and increase the dislocation opening stress.
Under the action of creep high temperature and stress, the degree of interface bonding between the reinforcement and the matrix will gradually decrease until debonding occurs, resulting in micropores; When the dislocation movement is hindered by the TiB whisker, dislocation plugging will occur, and the stress concentration will lead to cracking in the middle of TiB, and the standing and polymerization of microcracks will cause premature fracture of the alloy, which will reduce the creep life.
Microscopic defects are more generated in TiB whiskers, because the length-to-diameter ratio of TiC particles is small, and their approximate equiaxed shape effectively limits the stress concentration. The length-to-diameter ratio of long strips of TB whiskers is significantly larger than that of TiC particles, so stress concentration and fracture are more likely to occur.
The fracture morphology of the alloy after creep under 550°C and different stress conditions after descending two-stage aging treatment is shown in Fig. 12 and Fig. 13.
Fig. 12 Creep fracture of HT1 treated alloy (550°C)
There are large and small tenaments on the surface of the variable mouth, showing a typical ductile fracture pattern. The size of the large tendon socket is not only much larger than that of the small putty, but also the depth is significantly larger, indicating that the heat-treated alloy has high temperature creep toughness, corresponding to its larger creep strain (up to 92.8%).
At the same time, the disbonding of TiB and the cracking of TiB and TiC were widely observed near the ligament fossa, which once again confirmed the bearing effect of the mixed enhanced phase (TiB+TiC) on the external force during creep, which was consistent with the creep microstructure analysis results.
Fig. 13 Creep fracture of HT2 treated alloy (550°C)
conclusion
After the two-stage aging treatment of solid solution + descending order of the wrought alloy, fine dispersed secondary α phases were precipitated and more primary α phases with different forms were retained, and the grain boundary α phase was broken and spheroidized. The primary α phase is conducive to maintaining plasticity, but reduces the dispersion strengthening effect. After the solid solution + descending two-stage aging treatment at 860°C, the primary α phase disappeared, and the single secondary α phase structure caused the tensile strength of the alloy to increase greatly to 1734MPa, and the elongation decreased to 0.8%.
The descending two-stage aging treatment of the alloy has a time-varying life of more than 60h under the condition of 500C/300MPa. However, the creep curves at 550°C and different stresses showed typical three-stage characteristics of creep, and with the increase of stress at the same temperature, the steady-state creep rate of the alloy increased and the creep life decreased. The creep deformation mechanism of the alloy after the descending two-stage aging treatment is all dislocation climbing. For the same creep conditions, the creep life of the alloy after solid solution + descending two-stage aging treatment at 790°C is higher and the creep performance is better.
The secondary α phase deteriorates at variable high temperatures, and the equiaxed α phase is not conducive to improving the variable resistance due to the high diffusion rate. TiB and TiC strengthen the matrix during high-temperature creep. After heat treatment, the alloy has high creep toughness, which is manifested as the ductile fracture mode of concentrated load of TiB and TiC.
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