Solid phase transition of Cu precipitates in a low carbon TRIP assisted steel
Introduction
The evolution of crystal structure of Cu precipitates in steels has attracted significant in recent years [[1], [2], [3], [4], [5], [6], [7], [8], [9], [10], [11], [12], [13], [14], [15], [16], [17], [18], [19]]. The study of Cu precipitates in steels has both practical and theoretical significance because of the precipitation strengthening effect [[1], [2], [3], [4],[12], [13], [14]] and the complex crystal structure [[5], [6], [7], [8], [9], [10], [11], [12], [13], [14], [15], [16], [17], [18], [19]]. Cu precipitates in steels experience a long and complex change from nucleation to growth and finally to the stable FCC structure [[7], [8], [9], [10], [11],13]. This includes the processes of the nucleation of Cu atoms [[20], [21], [22], [23], [24], [25], [26]], martensite shear phase transformation [7,8,10,11,13], and the delamination of twins and periodic stacking faults [9,10,13]. Moreover, the phase transition behavior of Cu precipitates is different in different alloy systems [1,[7], [8], [9], [10], [11],13].
In Fe–Cu and Fe–Cu–Ni model steel, Othen et al. [6,7] studied the crystal structure of Cu precipitates at 550 °C for different aging time and proposed that the evolution sequence of crystal structure of Cu precipitates was BCC Cu(<4–5 nm)→9R Cu(4–17 nm)→3R Cu(18–30 nm)→FCC Cu(>40 nm) [7]. However, BCC Cu was speculated based on the structural characteristics of 9R Cu and was not directly observed. 3R Cu was proved by selected area electron diffraction (SAED) and it was actually FCC structure with a large number of defects. Monzen et al. [8,9] based on the research of Othen et al. [6,7], studied details of small size twin 9R Cu formed by direct phase transition of BCC Cu and found the elastic stress relaxation in them. They also proved a new crystal structure of Cu precipitates with FCT structure (a = b = 0.369 nm, c = 0.366 nm). A more detailed evolution sequence of Cu precipitation was BCC Cu (<5 nm at room temperature, < 12 nm 550 °C)→9R Cu (4–6 nm double twins, 7–15 nm multiple twins)→3R Cu (26–40 nm)→FCT Cu or FCC Cu (>40 nm). In Fe–3Si–2Cu-0.6Al-0.2Mn alloy system, Heo et al. [10] studied Cu precipitates nucleated on aging at different temperatures from 300 °C to 600 °C for 1 h. Nano BCC Cu clusters were directly observed in BCC Fe matrix. Non-twin 9R Cu, 9R Cu containing three twin variants (usually the number of 9R twin variants is an even number) and twinned FCC Cu were present in steels subjected to different aging temperatures. In addition, they proved that there was no 3R Cu, and the 9R structure could be directly transformed into twinned FCC Cu or FCC Cu through ±a/3[1 0 0]9R Cu Schockley-type partial dislocations. Heo proposed a new sequence of BCC Cu→9R Cu→twinned FCC Cu→FCC Cu. In the ultra-low C alloy system of Fe-0.005C-0.7Mn-2.0Cu-2.5Ni, the evolution of crystal structure of Cu precipitates was BCC Cu→9R Cu→twinned FCC or FCT Cu→FCC Cu [11]. In this alloy system, both twinned 9R Cu and non-twin 9R Cu were observed, which belong to monoclinic 9R structure. In Fe-2.5Cu-1.5Mn-4.0Ni-1.0-Al multicomponent ferritic steel, after aging at 500 °C for different time, Wen et al. [1] discovered a new sequence of B2 Cu–Ni–Mn–Al nano cluster (∼2 nm)→B2 Ni(Al,Mn) shell + Cu-rich BCC core (4–5 nm)→B2 Ni(Al,Mn) shell+9R Cu core (5–10 nm). Since Cu precipitates are surrounded by B2 Ni(Al,Mn) shells, their growth rate was very low. The strength of steel with a large number of B2 nanoclusters was more than double the strength of ordinary ferrite steels.
Quench-tempering heat treatment process is a conventional process for developing high strength low alloy matensite/bainite steel [[1], [2], [3], [4]]. Cu is also used as an alloy to acquire nano Cu precipitates in ferrite steels for precipitation-strengthening [[1], [2], [3]]. In order to simultaneously enhance plasticity and strength, regulation of residual austenite and nano precipitates in martensitic/bainite steels is a novel technology. Q&P process [[27], [28], [29]] has been used to develop ultra-high strength, higher ductility automotive sheet steel by obtain retained austenite in the martensitic matrix [30]. In order to obtain retained austenite in the thicker high strength low-alloy plate steel without adding a large number of alloy elements (Mn, Ni, C etc.), multi-step intercritical heat treatment process is developed [[31], [32], [33], [34], [35], [36], [37]] for enrichment the stabilizing elements in the reversed austenite. In our experimental low alloy steels, retained austenite and nano-precipitates were both obtained by two-step intercritical heat treatment. After intercritical tempering for different time, the process of reversing austenite comprising of nucleation, growth and stabilization was determined and the evolution behavior of Cu precipitates in the BCC Fe matrix with time is elucidated. A detailed study of Cu precipitates in high strength low alloy steels with retained austenite at different intercritical heat treatment time has been carried out [[12], [13], [14]]. In Fe-0.087C-2.05Mn-0.79Si-1.04Al-2.12Cu-1.50Ni-0.4(Mo+Nb) system, the phase transition sequence of Cu precipitates is Cu-rich nano-order cluster→9R Cu→detwinned 9R Cu→FCC Cu. The steel containing B2 FeCu nanoclusters showed a yield strength of 984 MPa with an increase of 226 MPa after a specific heat treatment [14].
In sequel to our previous study, the complex phase transition behavior of Cu precipitates in TRIP-assisted steels is elucidated. The solid phase transition of Cu precipitates from Cu-rich nano-order clusters to FCC Cu were studied by HR-TEM and 3DAP. A more complete phase transition process of Cu precipitates is revealed.
Section snippets
Experimental procedure
The experimental steel was a low-carbon low-alloy high-strength steel of nominal chemical composition Fe-0.087C-2.05Mn-0.79Si-1.04Al-2.12Cu-1.50Ni-0.4(Mo + Nb) in wt.%. The steel was subjected to two-step intercritical heat treatment. A good combination of strength and ductility with yield strength of 700–1000 MPa and total elongation of 20–30% were obtained through standard tensile test. The exact temperature of heat treatment, microstructure, mechanical properties and formation of reverted
The evolution of morphology, size, density, composition and volume fraction of Cu precipitates with intercritical tempering time
The evolution of morphology, size, density, composition and volume fraction of Cu precipitates in the BCC Fe matrix of steel on intercritical tempering from 1 min to 60 min was revealed by 3DAP results as shown in Fig. 1(a)-(l). Cu atoms began to separate from the matrix and nucleate at intercritical tempering time of 1 min. Fig. 1(a) shows the three-dimensional atom maps of Cu and 10 at.% iso-concentration surfaces of Cu-rich nano-clusters. Several small Cu precipitates in the range of
The evolution behavior of Cu precipitates in steel
Fig. 7 shows the evolution behavior of Cu precipitates in the low carbon low alloy steel when intercritically tempered at 680 °C. There are twenty-two high resolution images in Fig. 7, including Cu-rich nano clusters (three images), twinning 9R Cu (eight images), detwinned 9R Cu (three images), a complex transient Cu precipitation with both 9R and FCC structure (one image) and FCC Cu (four images), also non-twin 9R Cu (three images). The high-resolution images labeled I, II and III in Fig. 7
Conclusions
The evolution behavior of Cu precipitates during intercritical tempering in a low carbon low alloy steel was studied. A more complete crystal structure transformation process of Cu precipitates was revealed. The formation of Cu precipitates with different crystal structures and the transformation process of Cu precipitates was analyzed.
- 1.
The entire evolution process of Cu precipitates was caused by the continuous enrichment of Cu atoms in Cu precipitates and the size of Cu precipitates.
- 2.
The
Acknowledgements
The authors gratefully acknowledge the support from the National Natural Science Foundation of China (No. 51701012) and the Fundamental Research Funds for Central Universities (No. FRF-IC-18-009). R.D.K. Misra acknowledges support from Freeport McMoRan Fund 30203530 and continued with USTB.
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