Full length articleStrain hardening in Fe–16Mn–10Al–0.86C–5Ni high specific strength steel
Graphical abstract
Introduction
Steels are the strongest ductile bulk materials currently available [1], [2], [3], [4], [5], [6], [7]. High-strength and high-ductility steels are needed in various industrial sectors such as automobiles, aviation, aerospace, power, transport, and building construction. Such steels have been developed based on several design principles; typical categories include the transformation-induced plasticity (TRIP) steels [8], [9], twinning-induced plasticity (TWIP) steels [10], [11], dual-phase (DP) steels [12], [13], nano-structured steels [14], [15], and even hypereutectoid steel wires with ultrahigh (6.35 GPa) tensile strength [16].
Recently, the high-aluminum low-density steels have been actively studied for the purpose of increasing the specific strength (i.e. yield strength-to-mass density ratio) [4], [5], [17], [18], [19], [20], [21], [22], [23], [24], [25], [26], [27]. These low-density steels, mainly based on an Fe–Al–Mn–C alloy system containing high contents of Mn (16–28 wt.%), Al (3–12 wt.%) and C (0.7–1.2 wt.%), consist of face-centered cubic (fcc) austenite matrix and body-centered cubic (bcc) ferrite matrix and finely dispersed nanometer-sized κ-carbides of (Fe,Mn)3AlC type (the so-called TRIPLEX steel) [18], [19]. Recently, Kim et al. [27] showed a new variation: their Fe–16Mn–10Al–0.86C–5Ni high specific strength steel (HSSS) has a hard FeAl-type (B2) intermetallic compound as the strengthening second phase, and the alloying of Ni catalyzes the precipitation of B2 particles in the fcc matrix. The combination of specific strength and elongation is outstanding for this HSSS, when compared with other high-specific-strength alloys [27].
In developing various high strength steels aforementioned, a primary issue is the strain hardening capability. Strain hardening is a prerequisite for large uniform tensile ductility. However, the mechanism of strain hardening remains an open issue for most high strength steels [5], [22], [27], because they deform very heterogeneously due to their inhomogeneous microstructures. Even for an initial single-phase alloy, for instance TWIP and TRIP steels, deformation twins and martensite formation make the strain hardening behavior complex [20], [22]. Bhadeshia [28] pointed out that in TRIP steels it is unlikely that the large tensile elongation is predominantly caused by the transformation from austenite into martensite alone. The martensite colonies act as strong inclusions, akin to reinforcing components in a composite, and should have also played a role in strain hardening [29], [30]. A similar conclusion has been reached for TWIP steels where the twinning strain itself makes a significant though small contribution to the total elongation [31].
Indeed, plastic deformation in most high strength steels, much like in composites, is characterized by pronounced plastic heterogeneity between the constituent phases, as well as among grains with different orientations and mechanical responses towards an externally applied load. This causes complex internal stresses, which develop because of intra- and inter-granular variations of plastic strain. The load redistribution and strain partitioning resulting from the microstructural heterogeneity enable a high capacity of strain hardening, affecting the large ductility. The high internal stresses have in fact been reported to contribute to strain hardening and large ductility in TRIP steels [32], TWIP steels [33], [34], [35], nano-composites [36], [37], and dual-phase alloys [38], [39].
The development of internal stresses during deformation of an inhomogeneous microstructure with yield stress mismatch has been well described before [13], [40], [41]: upon tensile loading, plastic yield starts in the soft phase, and the applied load will be transferred from the soft phase to the hard one that is still in elastic state. Thus, internal stresses will build up at the phase interfaces. Upon unloading, the macroscopic stress remains higher than the stress in the soft phase until it reaches the unloaded state, where the soft phase is subjected to an (elastic) compression stress (a tensile stress in the hard phase) [42], [43]. If the two-phase alloy is subsequently subjected to compressive loading, it initially behaves elastically until the soft phase enters the plastic regime in compression, a situation that will take place at a much lower absolute stress compared to the tensile loading case because of the initial compression of the soft phase. A consequence is an asymmetry in the forward (tensile) and reverse (compressive) yield stresses. Such a phenomenon is known as the Bauschinger effect [35], [38], [44]. Recently, the use of diffraction techniques has supplanted this macroscopic description of internal stresses by the measurements of lattice strains in individual phases [45], [46]. The unload-reload tests [36], [38], [47] are also used for the study of internal stresses in thin films or composite wires where compression cannot be applied.
In this paper, we analyze the strain hardening in the Fe–16Mn–10Al–0.86C–5Ni HSSS composed of a γ-austenite matrix containing the B2 FeAl second phase. Based on in situ high energy X-ray diffraction data, the lattice strain evolution in both phases has been monitored and then used to correlate with the mechanical responses such as the stress and strain partitioning, the elasto-plastic transition and co-deformation, and the back-stress- induced strain hardening. Different from Kim et al. [27], who treated this HSSS as a case of precipitation strengthening with brittle and non-deformable B2 FeAl, here we show that this steel is better understood as a dual-phase microstructure, with plastic behavior much like a composite. In particular, the B2 phase is deformable, with significant strain hardening capability.
Section snippets
Materials
Similar to the procedures in Ref. [27], an Fe–16Mn–10Al–0.86C–5Ni (wt.%) HSSS was produced using arc melting in a high frequency induction furnace under pure argon atmosphere, and then cast to a cylindrical ingot with a diameter of 130 mm and length of 200 mm. The actual chemical composition of the ingot was determined to be Fe–16.4Mn–9.9Al–0.86C–4.8Ni–0.008P–0.004S (wt.%). The ingot was homogenized at 1180 °C for 2 h, hot forged in between 1150 °C and 900 °C into slabs with a thickness of
Microstructural characterization
Fig. 1a is an optical microscope image of the longitudinal section of the hot forged sample. Two phases are visible. One is the equi-axed grains of recrystallized fcc γ-austenite, while the other is the thick lamellar B2 phase parallel to the rolling direction. After cold rolling with a rolling strain of 80%, the γ grains change from granular to elongated shape, as shown in Fig. 1b. The B2 phase exhibits a much reduced thickness, indicative of its capability for plastic deformation. After
Plastic deformation in HSSS
Fig. 6 shows three stages of the lattice strain evolution with applied axial strains in the γ and B2 phases. In stage I, both γ and B2 phases deform elastically, with a linear increase of both and . In stage II, as shown in Fig. 6c, is the first to deviate from linearity at a strain of 0.06, while continues to rise linearly. This indicates the onset of yielding in γ, at the microscopic level, due to its lower yield stress than that of B2 [39]. Exactly from this very
Conclusions
We have analyzed the strain hardening process in the Fe–16Mn–10Al–0.86C–5Ni high specific strength steel. As the steel has a heterogeneous microstructure composed of a γ-austenite matrix containing the B2 FeAl second phase, our treatment is from the perspective of a dual-phase that are both deformable with significant strain hardening capability. In situ high-energy X-ray diffraction revealed the lattice strain evolution in both phases. The softer phase deformed first, shedding load to harder
Acknowledgments
This work was financially supported by the National Natural Science Foundation of China (NSFC) under grant Nos. 11572328, 11072243, 11222224, and 11472286, the 973 Programs under grant Nos. 2012CB932203, 2012CB937500, and 6138504. E.M. was supported at the Johns Hopkins University by U.S.-DOE-BES, Division of Materials Sciences and Engineering, under Contract No. DE-FG02-09ER46056.
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