High nitrogen nickel-free austenitic stainless steels (nitrogen concentration is usually higher than ∼ 0.4% in mass percent, termed as HNSs) have been studied extensively for their unique mechanical properties, i.e., improved strength and ductility, adequate work hardening ability, high fracture toughness and desirable resistance to corrosion, etc.^[1–11] These attractive properties depend on the benefit roles of nitrogen in enhancing solid-solution strengthening and stabilizing austenitic structure.^[1,3–6,8] The high potential of nitrogen to stabilize austenitic structure would make the expensive nickel be replaced and offer an additional advantage of cost-saving.^[1,3,4,6,8] Moreover, the enhanced solubility of nitrogen with the increase of the contents of alloying elements, i.e., manganese and chromium, can be advantageous for it to be effectively applied in the manganese and chromium-based stainless steels.^[1,3,8]

Extensive researches focused on the tensile mechanical properties and the microstructure evolution of the HNSs or other nitrogen alloyed austenitic stainless steels for extending their engineering applications have been reported.^[12–31] In these studies, the effects of various kinds of internal and external factors, such as grain size,^[12,16] twin structure (e.g., size, shape, and amount of twin),^[17,23–25] alloying elements (e.g., contents of carbon and nitrogen)^{[13–15,21,26,30,31]} and testing temperature and strain rate^{[18–20,22,27–29]} on the plastic deformation behaviors have been investigated extensively. Nevertheless, it is noted from these studies that very limited efforts were made to study the strain rate and cold rolling dependence of tensile strength and ductility of the HNS at a very wide range of strain rate and cold rolling deformation.

In this paper, the HNSs with a variety of initial microstructure and a very wide range of the strength and ductility were obtained through solution and cold rolling treatments. The corresponding tensile strength and ductility are investigated by performing tensile tests under a wide strain rate ranging from 10⁻⁴ s⁻¹ to 1 s⁻¹ at room temperature. The experimental tensile stress–strain curves were then examined by using the Ludwigson equation. Based on the experimental results of the stress–strain curve, the transmission electron microscopy observation of the microstructure and the scanning electron microscopy characterization of the deformation surfaces of this HNS, the significant strain rate and cold rolling dependence of the strength and ductility and the underlying plastic deformation mechanism was discussed and connected with the variation in the dislocation activity with strain rate and cold rolling.

The investigations on the metals and alloys have attracted extensive attention in the past years because of their good performance.^[32–39] For the purpose of optimizing the microstructure and mechanical properties of the metals and alloys and extending their practical applications, it is of particular importance to analyze their plastic deformation behavior or stress–strain curves. One of the earliest attempts is the empirical mathematical equation proposed by Hollomon (termed as the Hollomon equation).^[40] As well as the simplicity of the Hollomon equation, it can also explain the underlying deformation mechanisms and provide the correlation of the involved parameters with the microstructure features. Thus, this equation has been quite common to analyze the plastic deformation behavior of the polycrystalline metals during the monotonic tensile/compressive deformation over the decades. Based on the true stress-true plastic strain (σ–ε) curves, the Hollomon equation can be given by

In this study, the as-received HNS bar with a diameter of about 30 mm was first solution treated at 1150 °C for 8 hours followed by water quenching to form a uniform initial austenitic structure, and then it was cold rolled on a laboratory two-high mill at room temperature. The total reduction of area (RA) was about 10%, 30%, and 50%, respectively. The corresponding nominal chemical compositions have been reported in Ref. [51]. Uniaxial tensile tests (at least three times for each strain rate) were performed at room temperature on a MTS-810 test system with a broad strain rate ranging from 10⁻⁴ s⁻¹ to 1 s⁻¹. Dog-bone shaped specimens for tensile tests with a size of 8 mm × 2 mm × 1 mm (length × width × thickness) in the gauge section were cut from the HNS sheets with solution treatment and cold rolling condition by using an electrodischarging machine. Then these specimens were prepared by careful mechanically grinding and polishing to achieve a test surface with mirror-like finish.

Microstructure observation of the present HNS before testing was performed by field emission scanning electron microscopy (FESEM, SUPRA-40) with electron backscattering diffraction (EBSD) analysis. The obtained data for EBSD mapping were analyzed with a discrete step size in the range of 10–100 nm by using orientation imaging microscopy. Crystallographic structure of the present HNS before testing was examined by x-ray diffraction analysis (D/max 2500 PC) with Cu Ka radiation. Microstructure characterization of the present HNS before and after testing was carried out on a transmission electron microscopy (TEM, JEM-2100F) under an accelerating voltage of 200 kV respectively. To obtain the thin foil HNS for EBSD and TEM observation, the specimens cut from the bulk HNS were first mechanically ground to a thickness of ∼ 20 μm, then thinned by an ion beam with a tilt angle of 4°–8° using an Ion Polishing System (Gantan 691, USA) under a voltage of 5 kV. Morphologies of the deformation surfaces of the present HNS were examined by FESEM observation.

In polycrystalline metals, the statistically stored dislocations (SSDs) would generate initially in intragranular sources during the plastic deformation process. The glide of the SSDs is gradually in a form of planar or single slip and tends to distribute uniformly in the grain interiors.^[55,56] This glide of SSDs would undergo short-range resistance due to the local pinning of carbon and nitrogen interstitial atoms.^[49,54] As the deformation proceeds, the geometrically necessary dislocations (GNDs) would generate in the regions near the grain boundaries (GBs) to relax the stress/strain concentration.^[55,56] The glide of GNDs transforms gradually into the multiple slip on the intersecting slip systems^[55,56] and expands from the GB regions towards to the grain interiors with the increase of strain. The region of the multiple slip becomes larger and that of planar slip becomes smaller and disappears gradually. From different strain hardening roles of the planar slip and the multiple slip, the above dislocation activities can be clearly identified by the strain hardening exponents (n₁ and n₂) and the transition strain (ε_L) derived from the tensile stress–strain curves by using the Ludwigson equation. It has been suggested that n₁ is the measurement of the long-range stress associated with the interaction among the dislocations on the intersecting slip systems, n₂ is the measurement of the short-range stress due to the interactions of planar glide dislocations with solution interstitial atoms (e.g., carbon and nitrogen), and ε_L represents the strain at which a significant transition in the dislocation glide mode from the planer glide to the cross slip takes place. Therefore, the dependence of the dislocation activities on strain rate and cold rolling RA can be explained through the variations of n₁, n₂, and ε_L with strain rate and cold rolling RA.^[49,57,58]

For dislocations to nucleate and emit in intragranular sources, a critical (threshold) stress is required in the plastic deformation of polycrystalline metals. The subsequent plastic deformation would involve the multiplication and motion of dislocations and the interactions of dislocations on intersecting slip systems. The dislocation network created due to such dislocation activities would further hinder the motion of dislocations, which thus leads to the strain hardening. At high strain rate, owing to the high density or high storage rate of dislocations (SSDs and GNDs) generated initially in intragranular sources, a high initial critical stress or a high yield strength would be expected. As deformation proceeds, these high density or high storage rate of dislocations would enhance the multiple slip process. As a result, the transition in the dislocation glide mode from the planer glide to the multiple slip would occur at a smaller strain. Meanwhile, the enhanced interaction of dislocations would promote the dynamic recovery process of dislocation structure occur earlier or make the dislocation substructures form earlier (Figs. 8(b) and 8(e), Figs. 9(c) and 9(g)). Thus, a rapidly reduced strain hardening would occur at high strain rate. In contrast, a slowly reduced strain hardening would be expected at low strain rate. At large cold rolling RA, owing to the significant refinement of the grain or twin structure caused by cold rolling deformation, a high initial critical stress or a high yield strength would also be expected. As deformation proceeds, the reduced grain interior space would limit the initial planar slip process and make the transition in the dislocation glide mode from the planar glide to the multiple slip occurs at a much smaller strain. Meanwhile, the reduced grain interior space at large cold rolling RA would promote the dynamic recovery process of dislocation structure occur earlier or make the dislocation substructures form earlier (Figs. 8(b) and 9(c), Figs. 8(e) and 9(g)). Hence, a rapidly reduced strain hardening would occur at large cold rolling RA. In contrast, a slowly reduced strain hardening would take place at small cold rolling RA. Such variations in the strain hardening behavior are confirmed by the reduced n₁, ε_L, and n₂ with increasing strain rate and the increase of cold rolling RA (Table 1 or Fig. 6). It can also be connected with the strain rate and cold rolling dependence of the ultimate tensile strength of the present HNS. The increased ultimate tensile strength at high strain rate or large cold rolling RA arises from the higher yield strength in spite of the smaller stress increment caused by the rapidly reduced strain hardening (smaller n₁) (Fig. 4).

Owing to the inhomogeneous nature of plastic deformation of polycrystalline metals, a high strain hardening ability is generally needed to delay the deformation or strain localization process.^[59] At low strain rate, the development of the multiple slip and substructure would be more fully and uniformly over a wide strain range, which suppresses the deformation or strain localization process and delays the onset of plastic instability or necking until a larger strain level. Meanwhile, the grain deformation or GB sliding may also occur at low strain rate, which could delay the nucleation and growth of local microcracks by enhancing the deformation accommodation among the grains. At low strain rate, the strain hardening can continue over a larger strain range, a high tensile ductility (uniform and total elongations) is thus expected. In contrast, at high strain rate, the strain hardening would reduce rapidly as the strain increases. The nucleation and growth of local microcracks will occur at a low strain level, leading to a low tensile ductility at high strain rate. At small cold rolling RA, the development of the multiple slip and substructure would be more fully and uniformly due to the enough interior space, which suppresses the deformation or strain localization process and delays the onset of plastic instability or necking until a higher strain. At small cold rolling RA, the strain hardening can continue over a larger strain range, a high tensile ductility (uniform and total elongations) is thus expected. In contrast, at large cold rolling RA, the strain hardening would reduce rapidly as the strain increases. The nucleation and growth of local microcracks will occur at a low strain level, resulting in a low tensile ductility at large cold rolling RA. The above explanations for the strain rate and cold rolling dependence of the tensile ductility of the present HNS are well supported by the characterization results of the deformation surfaces shown in Fig. 10.

The tensile strength and ductility of a high nitrogen nickel-free austenitic stainless steel were investigated under tensile tests with a wide strain rate range at room temperature. The tensile tests revealed that this steel shows a strong strain rate and cold rolling dependence of the tensile strength and ductility. The analysis of the double logarithmic stress–strain curves indicated that this steel exhibits a two-stage strain hardening behavior, which can be examined and analyzed by using the Ludwigson equation. Based on the analysis results of the stress–strain curve, the transmission electron microscopy observation of the microstructure and the scanning electron microscopy characterization of the deformation surfaces, the strong strain rate and cold rolling dependence of the tensile strength and ductility of this steel was discussed and connected with the variation in the dislocation activities with strain rate and cold rolling. It reveals that the increased yield strength with the increase of strain rate and increasing cold rolling RA arises from the high density or high storage rate of dislocations generated at high strain rate and the reduced interior space caused by large cold rolling deformation, the increased ultimate tensile strength with the increase of strain rate and the increase of cold rolling RA arises mainly from the higher yielding before lower strain hardening ability (rapidly reduced strain hardening), the decreased uniform and total elongations with increasing strain rate and increasing cold rolling RA are attributed to the lower strain hardening ability (rapidly reduced strain hardening).