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Nanocrystalline metals are exceptionally strong but generally fail with only limited ductility when crack growth is controlled by dislocation emission from the crack tips. [1] A number of approaches were developed toward the production of nanocrystalline metals that can withstand significant plastic strain before failure, [2][3][4][5][6][7][8] including creating a bimodal grain size distribution, generating profuse nano-twins, forming nano-precipitates, and modifying the stacking fault energy. From the fracture viewpoint, these strategies are designed to introduce structural heterogeneities that limit crack propagation during plastic deformation.An alternative strategy that is not currently adopted in developing ductile nanocrystalline metals is based on the empirical correlation between ductility and the shear-to-bulk modulus ratio (G/B) which is inversely related to Poisson's ratio (n). [9,10] It is well known that isotropic microcrystalline metals with high n are generally more ductile by comparison to low-n metals of the same crystal structure [9] in spite of the exception for B2 intermetallic compounds such as NiAl that has a high n but is very brittle. [11] For an alloy system capable of forming bulk metallic glasses (BMGs), a brittle-to-ductile transition is also identified with increasing n and BMGs with n higher than % 0.31-0.32 usually exhibit high toughness. [10] Moreover, Pan et al. [12] showed that the volume of the shear transformation zones (STZs) in BMGs and ductility increase with increasing n, suggesting an intrinsic correlation between ductility and STZ volume. Nevertheless, no such correlation between ductility and n has been explored in nanocrystalline metals and alloys.On the fracture surface of a number of pure nanocrystalline metals, dimples are often observed with characteristic lengths of the order of % 100 nm to 2 mm which is considerably larger than the grain size but smaller than in coarse-grained polycrystals. [3][4][5] Similarly, fracture analysis on BMGs [13] reveals the formation of dimples (vein pattern) and the fracture toughness is proposed to depend on the length scale of the dimples that corresponds to the plastic process zone. Accordingly, the larger the plastic process zone, the tougher the BMG. It appears that nanocrystalline metals also behave mechanically in a manner similar to BMGs with a susceptibility to shear banding. [14][15][16] For example, shear bands COMMUNICATION [*] Dr.A nanocrystalline bcc Ti 67.4 Nb 24.6 Zr 5 Sn 3 alloy is shown to fracture in an intrinsically ductile manner with exceptionally large dimples (up to 10 mm) which are two orders of magnitude greater than the grain size (% 40 nm). This large plasticity length scale is attributed to a combination of low shear modulus (% 27 GPa), high Poisson's ratio (% 0.4) and ultrahigh strength (UTS % 1.1 GPa), close to the ideal shear stress, which facilitates ideal shear deformation to promote transgranular shear. 1108 wileyonlinelibrary.com ß
Nanocrystalline metals are exceptionally strong but generally fail with only limited ductility when crack growth is controlled by dislocation emission from the crack tips. [1] A number of approaches were developed toward the production of nanocrystalline metals that can withstand significant plastic strain before failure, [2][3][4][5][6][7][8] including creating a bimodal grain size distribution, generating profuse nano-twins, forming nano-precipitates, and modifying the stacking fault energy. From the fracture viewpoint, these strategies are designed to introduce structural heterogeneities that limit crack propagation during plastic deformation.An alternative strategy that is not currently adopted in developing ductile nanocrystalline metals is based on the empirical correlation between ductility and the shear-to-bulk modulus ratio (G/B) which is inversely related to Poisson's ratio (n). [9,10] It is well known that isotropic microcrystalline metals with high n are generally more ductile by comparison to low-n metals of the same crystal structure [9] in spite of the exception for B2 intermetallic compounds such as NiAl that has a high n but is very brittle. [11] For an alloy system capable of forming bulk metallic glasses (BMGs), a brittle-to-ductile transition is also identified with increasing n and BMGs with n higher than % 0.31-0.32 usually exhibit high toughness. [10] Moreover, Pan et al. [12] showed that the volume of the shear transformation zones (STZs) in BMGs and ductility increase with increasing n, suggesting an intrinsic correlation between ductility and STZ volume. Nevertheless, no such correlation between ductility and n has been explored in nanocrystalline metals and alloys.On the fracture surface of a number of pure nanocrystalline metals, dimples are often observed with characteristic lengths of the order of % 100 nm to 2 mm which is considerably larger than the grain size but smaller than in coarse-grained polycrystals. [3][4][5] Similarly, fracture analysis on BMGs [13] reveals the formation of dimples (vein pattern) and the fracture toughness is proposed to depend on the length scale of the dimples that corresponds to the plastic process zone. Accordingly, the larger the plastic process zone, the tougher the BMG. It appears that nanocrystalline metals also behave mechanically in a manner similar to BMGs with a susceptibility to shear banding. [14][15][16] For example, shear bands COMMUNICATION [*] Dr.A nanocrystalline bcc Ti 67.4 Nb 24.6 Zr 5 Sn 3 alloy is shown to fracture in an intrinsically ductile manner with exceptionally large dimples (up to 10 mm) which are two orders of magnitude greater than the grain size (% 40 nm). This large plasticity length scale is attributed to a combination of low shear modulus (% 27 GPa), high Poisson's ratio (% 0.4) and ultrahigh strength (UTS % 1.1 GPa), close to the ideal shear stress, which facilitates ideal shear deformation to promote transgranular shear. 1108 wileyonlinelibrary.com ß
The present work is aimed at characterizing the strain hardening behavior of AISI 316L austenitic stainless steel using ultrasonic velocity measurements. For this purpose, microstructural studies and ultrasonic velocity measurements were carried out on the samples deformed to different levels of strain at room temperature. Strikingly, the ultrasonic velocity-strain plot of the alloy exhibited a three-stage behavior that was similar to the strain hardening rate-strain response of the alloy. At strains lower than about 0.06 (stage A), a falling regime of velocity was observed that was related to the increase of dislocations density. This stage was followed by a stage of a nearly constant velocity (stage B). The initiation of this stage was concurrent with the onset of deformation twinning in the microstructure. Beyond a strain of about 0.2, the second falling regime of velocity (stage C) was developed. The occurrence of this stage was attributed to the difficulty of new twins formation with an increasing strain.In recent years, ultrasonic velocity measurement has been used increasingly as a nondestructive technique to study various characteristics and processes of materials. Several investigations have been performed to correlate velocity measurements with the microstructural features, [1][2][3] physical properties, [4] mechanical properties, [5][6][7] heat treatments, [8][9][10] grain size, [11] and the amount of deformation. [12][13][14] However, the correlation between strain hardening behavior and ultrasonic velocity, especially in low stacking fault energy (SFE) alloys, has not received attention so far.The strain hardening technique is widely used for strengthening the metallic components. From a technological viewpoint, an accurate characterization of strain hardening behavior of a material is necessary to develop models with good predictive capabilities concerning the force-displacement relationship in deformation processes. In addition, it provides clear indications of the beginning or collapse of many of the microstructural phenomena occurring during deformation. A transmission electron microscope (TEM) and an optical microscope (OM) usually are applied to characterize the strain hardening behavior of materials; however, sample preparation for these techniques is difficult and time consuming. Also, in a recent study, it has been shown that the deformation mechanism of SFE alloys cannot be characterized solely by OM. [15] Strain hardening behavior of low SFE face-centeredcubic (fcc) alloys (such as Hadfield steels, austenitic stainless steels, and multiple Co-Ni superalloys) have been studied extensively. [16][17][18][19] It is known that the strain hardening rate (h)-true strain (e) plots of these alloys exhibit four distinct regimes of hardening. In stage A, the strain hardening rate decreases continuously with deformation. This stage is followed by a second stage of nearly constant hardening rate, stage B. The second falling regime of strain hardening is observed in stage C. At large strains, this stage...
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