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Tensile behavior of nanocrystalline Ni–Fe alloy
Materials Science and Engineering A363 (2003) 62–66
Tensile behavior of nanocrystalline Ni–Fe alloy X.Y. Qin a,∗ , S.H. Cheong b , J.S. Lee b a
Key Lab. of Internal Friction and defects in Solids, Institute of Solid State Physics, Academia Sinica, 230031 Hefei, PR China b Department of Metallurgy and Material Science, Hanyang University, 425-791 Ansan, Republic of Korea Received 17 March 2003; received in revised form 6 June 2003
Abstract Tensile deformation and fracture behavior of bulk nanocrystalline Ni–Fe alloy, synthesized by a mechano-chemical method plus hot-isostatic pressing, were investigated by using mechanical tests and fractography. The results indicate that the as-prepared specimen fails prematurely with fracture strength (σ f ∼1.15 GPa) lower than yield stress evaluated from hardness measurements. However, σ f increases monotonously with time in isothermal (850 ◦ C) annealing, in contrast to which microhardness (yield strength) decreases with the annealing time. Fractography indicates that for the as-prepared specimen inter-particle brittle cracking is the main fracture mode. Nevertheless, microscopically local ductile fracture morphology was revealed to exist as manifested by the appearance of nano-dimples. When annealing time ta > ∼1.2 h σ f surpasses yield strength; correspondingly fractures exhibit more ductile mode, with a large number of both (sub-) micrometer-scale and mesoscale dimples prevailing as ta reaches 24 h. The improvement of mechanical properties and the change of fracture morphology with annealing are discussed. © 2003 Elsevier B.V. All rights reserved. Keywords: Nanocrystalline material; Alloy (Ni–Fe); Tensile deformation and fracture
1. Introduction The mechanical behavior and plastic deformation mechanisms of nanocrystalline materials (n-materials) have been studied experimentally using various methods which include normal hardness measurements [1], indentation [2], creep [3], bending tests [2,4], compression [5], tensile tests [6] and so on. Although some phenomena were revealed, such as high strength [5,6], low tensile strain [6] and perfect plastic behavior [2,5], the deformation and fracture mechanism in these materials are still a controversy issue. Further exploration on this problem is clearly vitally significant from viewpoints of fundamental research as well as applications. Here, we present our investigations on tensile deformation and fracture characteristics of bulk nanocrystalline ␥-Ni–19Fe alloy (n-Ni–Fe), synthesized by a mechano-chemical process plus hot isostatic pressing (HIPing). ␥-Ni–xFe with x = 10–65 wt.%, so called Permalloy, are important materials with wide applications in industry [7]. Many techniques [8–13] have been utilized to prepare ∗ Corresponding author. Tel.: +86-55-15-592750; fax: +86-55-15-591434. E-mail address: [email protected] (X.Y. Qin).
0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0921-5093(03)00497-0
n-Ni–Fe, mostly in powder state. However, mechanical properties of bulk n-Ni–Fe with nearly full density were rarely reported except hardness tests [8]. Recently, we reported the compressive behavior of the n-Ni–Fe [14,15]. In order to understand further its mechanical properties, tensile deformation and fracture experiments for n-Fe–Ni were conducted. Emphasis was laid on combined investigations of mechanical tests with fracture morphology.
2. Material syntheses and experimental procedures 2.1. Synthesis and characterization of the material The powders of n-Ni–Fe alloy were synthesized by using a mechano-chemical process [13,16]. Briefly, ␣-Fe2 O3 (30 m, 99.99%) and NiO (4 m, 99.99%) powders were blended and ball milled. Then they were reduced at 500 ◦ C for 1 h and alloyed at 550 ◦ C for 0.5 h in hydrogen. The obtained Ni–Fe powder (∼20 nm) were compacted and pre-sintered at 650 ◦ C for 1.5 h. Highly dense bars were prepared with HIPing at 750 ◦ C for 1 h under the 190 MPa in argon. The obtained Ni–Fe alloy has ␥ phase (fcc structure). The mean grain sizes were determined from XRD
X.Y. Qin et al. / Materials Science and Engineering A363 (2003) 62–66
to be 37 nm. The bulk density was measured based on Archimedes principle to be 98.5% theoretical one. The concentrations (wt.%) of the alloy are Fe19.4, Cr0.15, Mn0.05, Al0.05, and Si0.05%.
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2.2. Mechanical tests and fractography Dog-bone shaped tensile specimens with the gage length of 6 mm and gage cross section of 2 ×(∼1) mm2 were cut from HIPed bars (typical size: ∼23 ×(∼7) ×(∼1.2) mm3 ) with a spark erosion machine. To avoid surface damage, a layer of more than ∼0.1 mm in thickness was polished away from all sides (especially in the gage section) first with grinding papers and then with diamond paste. To investigate annealing effects on its tensile behavior, as-HIPed specimens were annealed isothermally at 850 ◦ C for different time in pure (5 N) argon. Tensile tests were carried out at ambient temperature on an Instron (model: 1195) machine with a strain rate of ∼1.4 ×10−4 s−1 . Strain was obtained by recording time multiplied by crosshead speed. Because of brittle nature of the materials, yield strength could not be determined accurately from the stress–strain curve due to premature failure, and therefore Vickers microhardness was measured to evaluate its yield strength, where a load of 100 g was applied and a residence time of 20 s was used. Fractography was carried out with optical microscope (OM) and field emission scanning microscope (FE-SEM).
3. Results 3.1. Mechanical tests Fig. 1 gives the stress-strain curves for the n-Ni–Fe specimens in as-HIPed state and annealed states (for clarity, the
Fig. 2. The variation of fracture stress and microhardness (divided by three: Hv /3) with time in isothermal annealing (at 850 ◦ C).
data of only two annealed (1.5, 24 h) specimens are plotted). It can be seen that the specimen in as-HIPed state shows very brittle behavior and fails in elastic region. After annealing at 850 ◦ C, its fracture stress and strain increases. However, no obvious yielding was observed for the annealed specimens except for that annealed for 24 h, whose strain–stress curve deviates from linear elastic regime at ∼1100 MPa. Due to this brittleness of n-Ni–Fe in tension, Vickers microhardness was determined to evaluate its yield behavior. Fig. 2 shows the microhardness divided by three as a function of annealing time ta . For comparison, its fracture strength σ f is also plotted vs annealing time. It can be seen that the value of the hardness divided by three (Hv /3) is much higher than fracture stress σ f for the specimen in as-HIPed state. Then Hv /3 drops promptly from ∼1.48 GPa in as-HIPed state down to 1.07 GPa in the state annealed for ta = 6 h. As annealing time is prolonged further, the hardness continues to decrease but with a much small decreasing rate. On the other hand, with increase of ta fracture strength σ f increases from 1.15 GPa (in as-HIPed state) monotonously to ∼1.38 GPa as ta = 24 h, intersecting with the plot of Hv /3 vs ta at ta = ∼1.2 h. Since a relationship σ y = Hv /3 between yield strength σ y and Vickers hardness holds approximately in n-Ni–Fe [16], the fact that fracture strength σ f > Hv /3 (≈σ y ) as ta > ∼1.2 h indicates that some macroscopic plastic deformation occurs before failure in the specimens annealed for ta > ∼1.2 h. The larger difference between σ f and Hv /3 after further annealing suggests that much more plastic deformation would occur before failure and fracture mode would become more ductile in nature for those annealed for longer time. 3.2. Fractography
Fig. 1. Engineering stress–strain curves for an as-HIPed specimen (䊐) and the specimens annealed at 850 ◦ C for 1.5 (䊉) and 24 h (䊊).
OM observations revealed no detectable plastic deformation traces near fractures for the as-HIPed specimen (not shown here). In contrast, some plastic traces (necking) were observed for annealed specimens (especially for the
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Fig. 3. Fracture surface morphologies observed with FE-SEM for the as-HIPed specimen: (a) at the magnification of ×500; (b) at the magnification of ×5000 (the arrow indicates one of enlarged vein-like patterns as shown in (a)); (c) at the magnification of ×50,000.
specimen annealed for 24 h). However, for all the specimens tested, regardless of the fracture mode (brittle or more ductile), the macroscopic fracture surfaces were perpendicular to loading axis, although locally fractures deviate from their main directions with irregular shape left. Fig. 3 shows FE-SEM fractographs for the as-HIPed specimen with various magnifications. At low magnifications (< ×300), the fracture surface looks very smooth and featureless; however, at higher magnifications, ×500 for instance, the surface displays some vein-like patterns (Fig. 3a) that are similar to what was frequently observed in metallic glasses. At the magnification of ×5000 (Fig. 3b), one can see clearly that the fracture surface displays granular structure, with a typical granular size being in sub-micrometer range. Since the mean grain size for this specimen is around 40 nm, the granular structures with the characteristic dimension of sub-micrometer scale as shown in Fig. 3b are in fact corresponding to agglomerate particles before HIPing [13]. This result suggests that the crack initiates and
propagates along interfaces between the particles leading to brittle inter-particle fracture. The vein-like patterns observed at low magnification (Fig. 3a) can be seen more clearly in Fig. 3b (as shown by an arrow). It can be seen from this figure that this pattern is composed of ridges that form seemingly through the meeting of the cracks propagating in different (but nearly parallel) planes and tearing off. Hence, intra-particle cracking would have happened in these regions. Fig. 3c gives a micrograph of the region with the vein-like pattern at a much high magnification. From this graph one can see that some nano-grains are recognizable. However, the interesting thing one can find in this graph is the appearance of nano-scale dimples with typical size ranging from ∼50 to ∼200 nm. This result indicates that although n-Ni–Fe shows brittle behavior macroscopically, microscopically ductile fracture morphology can still exist locally. The fracture behavior of n-Ni–Fe changes to more ductile mode after annealing, and with the increase of annealing time this change become much clear. Fig. 4 shows the
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4. Discussion
Fig. 4. Fracture surface morphologies observed with FE-SEM for the specimen annealed at 850 ◦ C for 24 h: (a) at the magnification of ×5000; (b) at the magnification of ×50,000.
FE-SEM micrographs of the fractures for the specimen annealed for 24 h. At low magnifications (×500, for instance) its fracture morphology (not shown here) is similar to that (Fig. 3a) for as-HIPed specimens, except for appearance of some porous like structure. However, its morphology at higher magnification (×5000) (Fig. 4a) is clearly different from that for as-HIPed specimen (Fig. 3b). One can see that the fracture surface is filled with dimples with their sizes being in the range of 0.2–2 m, as distinguished from the granular feature of the as-HIPed specimen (Fig. 3b). Fig. 4b gives the fractograph at much high magnification for the same fracture surface. This graph reveals that microscopically the annealed specimen displays very ductile fracture mode: mesoscale dimples appear with grains locating frequently inside. However, grain growth takes place substantially in this specimen, with some of them growing beyond 100 nm, although there are still some grains with sizes below 50 nm.
Tensile tests show that as-HIPed n-Ni–Fe fails prematurely in elastic region and its ultimate strength is only ∼1.1 GPa. This value is much smaller than the yield strength (∼1.6 GPa) obtained from compression tests [15] and that (∼1.5 GPa) evaluated from hardness measurements here, indicating asymmetry between tension and compression. This mechanical behavior is presumably related to the defects present in the specimens. Since our material is synthesized through powder fabrication and subsequent compaction by using HIPing, interfaces between agglomerate particles formed during HIPing should be weak places, where some microflaws or unclosed pores may exist due to residual unformed bonding. This speculation is reasonable, considering that the relative densities of the specimens are less than 100% of the theoretical one. On the other hand, deformation through dislocation activity in this material is expected to be difficult due to its small grain size (37 nm), considering that multiplication and slip of dislocations within these small grains need too much high stress. Furthermore, grain boundary sliding at room temperature can hardly happen in this material [15]. Therefore, under action of applied external stresses cracking would first initiate in the interfaces or existing microflaws there would propagate before any plastic deformation occurs, leading to premature failure of the material. The explanation of premature failure through interfacial cracking is consistent with the fractography, where the granular structures with the characteristic dimensions of agglomerate particle sizes appear (Fig. 3b). The observed vein-like patterns at low magnification (Fig. 3a) for as-HIPed specimen are seemingly similar to that observed in metallic glasses. But, their microscopic details are different. Vein patterns observed in metallic glass are usually featureless, and have no microstructural details [17]; while vein-like patterns observed here have, as mentioned above, nano-scale dimples. Since these vein-like patterns are presumably caused through tearing off, high rates of stain and cracking propagation along interfaces between particles are expected, which would give rise to local temperature rise. This would facilitate coalescence of free volumes, and so lead to formation of the nano-dimples. After the specimens are annealed at higher temperatures, fracture strength increases and hardness (yield stress) decreases, corresponding to which the fracture mode changes to more ductile manner. This change behavior should result from its microstructural and/or defect evolution during annealing. This change of microstructures may include: (1) grain growth; (2) enhancement of inter-particle adhesion; (3) void (porosity) formation coming from coalescence of residual micro-voids or free volumes. It is evident that grain growth (comparing Fig. 3c with Fig. 4b) may soften the material and enable the intra-granular deformation to become easier, leading to decrease in its yield stress (hardness) with annealing time (see Fig. 2). While inter-particle adhesion enhancement should inhibit inter-particle cracking
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and postpone occurrence of interfacial failure, which lead inevitably to increase in fracture strength with annealing, as observed in mechanical tests (Fig. 2). The large ((sub-) micrometer-sized) dimples observed (Fig. 4a) in the fractures of annealed specimens can be ascribed to the voids (porosity) formed through coalescence of residual micro-voids or free volumes. However, we note that ductile fracture morphology in the annealed specimens could be divided into two different levels: micrometer scale and mesoscale. Apart from the large dimples (Fig. 4a) being related to void formation, on much finer scale the formation of a large number of mesoscale dimples with sizes of several hundred nanometers (Fig. 4b) could be related to plastic deformation due to their configurations and dense population. They would initiate through stress concentrations produced at smaller nano-grains and develop through intragranular deformation in larger grains, for nano-grains are frequently observed to locate inside mesoscale-dimples (Fig. 4b). It is understandable that smaller grains with higher resistance to dislocation motion should act as obstacles to deformation. Plastic deformation would therefore proceed in large grains around the smaller grains. In other words, the mesoscale-dimples would presumably be associated with both concentrating effect of smaller grains and increasing deformation ability in larger grains. This speculation is supported by the fact that as annealing time reaches 24 h yield strength (the hardness Hv /3) is much smaller than fracture strength (Fig. 2) and macroscopic plastic deformation occurs substantially (Fig. 1). 5. Conclusions The characteristics of tensile deformation and fracture for as-prepared and annealed n-Ni–Fe alloy were investigated. The results indicate that the as-prepared specimen fails prematurely with ultimate strength (σ f ∼1.15 GPa) lower than yield strength evaluated from its Vickers microhardness. However, σ f increases and the hardness (yield strength) decreases monotonously with annealing time. Fractography indicates that for the as-prepared specimen inter-particle cracking is the main fracture mode, with nano-dimples appearing locally. When annealing time ta > ∼1.2 h σ f surpasses yield strength; correspondingly fractures are filled with a large number of both (sub-) micrometer-scale and mesoscale dimples as ta reaches 24 h. The improvement of fracture strength upon annealing can be ascribed mainly to
the enhancement of interfacial strength, while the decrease in yield strength (hardness) with annealing would originate from grain growth.
Acknowledgements The author (Qin) gratefully acknowledges Alexander von Humboldt Foundation for providing a fellowship. Specially he wishes to acknowledge Professor E. Nembach of Institut für Metallforschung, Universität Münster who contributed much to this work through arrangement of facilities and effective instructions. The authors are indebted to T. Krol for his assistance in mechanical tests and Bodycote IMT GmbH, Essen, Germany for HIPing experiments. Follow-up support from Academic Sinica through ‘Hundred Person Program’ and Korean Ministry of Science and Technology through the ‘2001 National Research Laboratory Program’ is gratefully acknowledged.
References [1] A.M. El-Sherik, U. Erb, G. Palumbo, K.T. Aust, Scripta Metall. 27 (1992) 1185. [2] T.R. Malow, C.C. Koch, Metall. Mater. Trans. 29A (1998) 2285. [3] N. Wang, Z. Wang, K.T. Aust, U. Erb, Mater. Sci. Eng. A 237 (1997) 150. [4] V.Y. Gertsman, M. Hoffmann, H. Gleiter, R. Birringer, Acta Metall. 42 (1994) 3539. [5] R.Z. Valiev, B.Y. Kozlov, Y.F. Ivannov, J. Lian, A.A. Nazarov, B. Baudelet, Acta Metall. 42 (1994) 2467. [6] P.G. Sanders, J.A. Eastman, J.R. Weertman, Acta Metall. 445 (1997) 4019. [7] E.P. Wohlfarth (Ed.), Ferromagnetic Materials, vol. 2, North-Holland Pub. Co, Amsterdam, 1980, p. 123. [8] C. Cheung, F. Djuanda, U. Erb, G. Palumbo, Nanostruct. Mater. 5 (1995) 513. [9] C. Kuhrt, L. Schultz, J. Appl. Phys. 73 (1993) 1975. [10] F.D. Ge, L.M. Chen, W.J. Ku, J. Zhu, Nanostruct. Mater. 8 (1997) 703. [11] S. Eroglu, S.C. Zhang, G.L. Messing, J. Mater. Res. 11 (1996) 1231. [12] Y.H. Zhou, M. Harmelin, J. Bigot, Scripta Metall. 23 (1989) 1391. [13] X.Y. Qin, J.S. Lee, J.G. Nam, B.S. Kim, Nanostruct. Mater. 11 (1999) 383. [14] X.Y. Qin, X.G. Zhu, S. Gao, L.F. Chi, J.S. Lee, Scripta Mater. 46 (2002) 611. [15] X.Y. Qin, X.G. Zhu, S. Gao, L.F. Chi, J.S. Lee, J. Phys. Condens. Matter 14 (2002) 2606. [16] X.Y. Qin, J.S. Lee, C.S. Lee, J. Mater. Res. 17 (2002) 991. [17] Z. Bian, G. He, G.L. Chen, Scripta Mater. 46 (2002) 407.
Tensile behavior of nanocrystalline Ni–Fe alloy X.Y. Qin a,∗ , S.H. Cheong b , J.S. Lee b a
Key Lab. of Internal Friction and defects in Solids, Institute of Solid State Physics, Academia Sinica, 230031 Hefei, PR China b Department of Metallurgy and Material Science, Hanyang University, 425-791 Ansan, Republic of Korea Received 17 March 2003; received in revised form 6 June 2003
Abstract Tensile deformation and fracture behavior of bulk nanocrystalline Ni–Fe alloy, synthesized by a mechano-chemical method plus hot-isostatic pressing, were investigated by using mechanical tests and fractography. The results indicate that the as-prepared specimen fails prematurely with fracture strength (σ f ∼1.15 GPa) lower than yield stress evaluated from hardness measurements. However, σ f increases monotonously with time in isothermal (850 ◦ C) annealing, in contrast to which microhardness (yield strength) decreases with the annealing time. Fractography indicates that for the as-prepared specimen inter-particle brittle cracking is the main fracture mode. Nevertheless, microscopically local ductile fracture morphology was revealed to exist as manifested by the appearance of nano-dimples. When annealing time ta > ∼1.2 h σ f surpasses yield strength; correspondingly fractures exhibit more ductile mode, with a large number of both (sub-) micrometer-scale and mesoscale dimples prevailing as ta reaches 24 h. The improvement of mechanical properties and the change of fracture morphology with annealing are discussed. © 2003 Elsevier B.V. All rights reserved. Keywords: Nanocrystalline material; Alloy (Ni–Fe); Tensile deformation and fracture
1. Introduction The mechanical behavior and plastic deformation mechanisms of nanocrystalline materials (n-materials) have been studied experimentally using various methods which include normal hardness measurements [1], indentation [2], creep [3], bending tests [2,4], compression [5], tensile tests [6] and so on. Although some phenomena were revealed, such as high strength [5,6], low tensile strain [6] and perfect plastic behavior [2,5], the deformation and fracture mechanism in these materials are still a controversy issue. Further exploration on this problem is clearly vitally significant from viewpoints of fundamental research as well as applications. Here, we present our investigations on tensile deformation and fracture characteristics of bulk nanocrystalline ␥-Ni–19Fe alloy (n-Ni–Fe), synthesized by a mechano-chemical process plus hot isostatic pressing (HIPing). ␥-Ni–xFe with x = 10–65 wt.%, so called Permalloy, are important materials with wide applications in industry [7]. Many techniques [8–13] have been utilized to prepare ∗ Corresponding author. Tel.: +86-55-15-592750; fax: +86-55-15-591434. E-mail address: [email protected] (X.Y. Qin).
0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0921-5093(03)00497-0
n-Ni–Fe, mostly in powder state. However, mechanical properties of bulk n-Ni–Fe with nearly full density were rarely reported except hardness tests [8]. Recently, we reported the compressive behavior of the n-Ni–Fe [14,15]. In order to understand further its mechanical properties, tensile deformation and fracture experiments for n-Fe–Ni were conducted. Emphasis was laid on combined investigations of mechanical tests with fracture morphology.
2. Material syntheses and experimental procedures 2.1. Synthesis and characterization of the material The powders of n-Ni–Fe alloy were synthesized by using a mechano-chemical process [13,16]. Briefly, ␣-Fe2 O3 (30 m, 99.99%) and NiO (4 m, 99.99%) powders were blended and ball milled. Then they were reduced at 500 ◦ C for 1 h and alloyed at 550 ◦ C for 0.5 h in hydrogen. The obtained Ni–Fe powder (∼20 nm) were compacted and pre-sintered at 650 ◦ C for 1.5 h. Highly dense bars were prepared with HIPing at 750 ◦ C for 1 h under the 190 MPa in argon. The obtained Ni–Fe alloy has ␥ phase (fcc structure). The mean grain sizes were determined from XRD
X.Y. Qin et al. / Materials Science and Engineering A363 (2003) 62–66
to be 37 nm. The bulk density was measured based on Archimedes principle to be 98.5% theoretical one. The concentrations (wt.%) of the alloy are Fe19.4, Cr0.15, Mn0.05, Al0.05, and Si0.05%.
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2.2. Mechanical tests and fractography Dog-bone shaped tensile specimens with the gage length of 6 mm and gage cross section of 2 ×(∼1) mm2 were cut from HIPed bars (typical size: ∼23 ×(∼7) ×(∼1.2) mm3 ) with a spark erosion machine. To avoid surface damage, a layer of more than ∼0.1 mm in thickness was polished away from all sides (especially in the gage section) first with grinding papers and then with diamond paste. To investigate annealing effects on its tensile behavior, as-HIPed specimens were annealed isothermally at 850 ◦ C for different time in pure (5 N) argon. Tensile tests were carried out at ambient temperature on an Instron (model: 1195) machine with a strain rate of ∼1.4 ×10−4 s−1 . Strain was obtained by recording time multiplied by crosshead speed. Because of brittle nature of the materials, yield strength could not be determined accurately from the stress–strain curve due to premature failure, and therefore Vickers microhardness was measured to evaluate its yield strength, where a load of 100 g was applied and a residence time of 20 s was used. Fractography was carried out with optical microscope (OM) and field emission scanning microscope (FE-SEM).
3. Results 3.1. Mechanical tests Fig. 1 gives the stress-strain curves for the n-Ni–Fe specimens in as-HIPed state and annealed states (for clarity, the
Fig. 2. The variation of fracture stress and microhardness (divided by three: Hv /3) with time in isothermal annealing (at 850 ◦ C).
data of only two annealed (1.5, 24 h) specimens are plotted). It can be seen that the specimen in as-HIPed state shows very brittle behavior and fails in elastic region. After annealing at 850 ◦ C, its fracture stress and strain increases. However, no obvious yielding was observed for the annealed specimens except for that annealed for 24 h, whose strain–stress curve deviates from linear elastic regime at ∼1100 MPa. Due to this brittleness of n-Ni–Fe in tension, Vickers microhardness was determined to evaluate its yield behavior. Fig. 2 shows the microhardness divided by three as a function of annealing time ta . For comparison, its fracture strength σ f is also plotted vs annealing time. It can be seen that the value of the hardness divided by three (Hv /3) is much higher than fracture stress σ f for the specimen in as-HIPed state. Then Hv /3 drops promptly from ∼1.48 GPa in as-HIPed state down to 1.07 GPa in the state annealed for ta = 6 h. As annealing time is prolonged further, the hardness continues to decrease but with a much small decreasing rate. On the other hand, with increase of ta fracture strength σ f increases from 1.15 GPa (in as-HIPed state) monotonously to ∼1.38 GPa as ta = 24 h, intersecting with the plot of Hv /3 vs ta at ta = ∼1.2 h. Since a relationship σ y = Hv /3 between yield strength σ y and Vickers hardness holds approximately in n-Ni–Fe [16], the fact that fracture strength σ f > Hv /3 (≈σ y ) as ta > ∼1.2 h indicates that some macroscopic plastic deformation occurs before failure in the specimens annealed for ta > ∼1.2 h. The larger difference between σ f and Hv /3 after further annealing suggests that much more plastic deformation would occur before failure and fracture mode would become more ductile in nature for those annealed for longer time. 3.2. Fractography
Fig. 1. Engineering stress–strain curves for an as-HIPed specimen (䊐) and the specimens annealed at 850 ◦ C for 1.5 (䊉) and 24 h (䊊).
OM observations revealed no detectable plastic deformation traces near fractures for the as-HIPed specimen (not shown here). In contrast, some plastic traces (necking) were observed for annealed specimens (especially for the
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Fig. 3. Fracture surface morphologies observed with FE-SEM for the as-HIPed specimen: (a) at the magnification of ×500; (b) at the magnification of ×5000 (the arrow indicates one of enlarged vein-like patterns as shown in (a)); (c) at the magnification of ×50,000.
specimen annealed for 24 h). However, for all the specimens tested, regardless of the fracture mode (brittle or more ductile), the macroscopic fracture surfaces were perpendicular to loading axis, although locally fractures deviate from their main directions with irregular shape left. Fig. 3 shows FE-SEM fractographs for the as-HIPed specimen with various magnifications. At low magnifications (< ×300), the fracture surface looks very smooth and featureless; however, at higher magnifications, ×500 for instance, the surface displays some vein-like patterns (Fig. 3a) that are similar to what was frequently observed in metallic glasses. At the magnification of ×5000 (Fig. 3b), one can see clearly that the fracture surface displays granular structure, with a typical granular size being in sub-micrometer range. Since the mean grain size for this specimen is around 40 nm, the granular structures with the characteristic dimension of sub-micrometer scale as shown in Fig. 3b are in fact corresponding to agglomerate particles before HIPing [13]. This result suggests that the crack initiates and
propagates along interfaces between the particles leading to brittle inter-particle fracture. The vein-like patterns observed at low magnification (Fig. 3a) can be seen more clearly in Fig. 3b (as shown by an arrow). It can be seen from this figure that this pattern is composed of ridges that form seemingly through the meeting of the cracks propagating in different (but nearly parallel) planes and tearing off. Hence, intra-particle cracking would have happened in these regions. Fig. 3c gives a micrograph of the region with the vein-like pattern at a much high magnification. From this graph one can see that some nano-grains are recognizable. However, the interesting thing one can find in this graph is the appearance of nano-scale dimples with typical size ranging from ∼50 to ∼200 nm. This result indicates that although n-Ni–Fe shows brittle behavior macroscopically, microscopically ductile fracture morphology can still exist locally. The fracture behavior of n-Ni–Fe changes to more ductile mode after annealing, and with the increase of annealing time this change become much clear. Fig. 4 shows the
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Fig. 4. Fracture surface morphologies observed with FE-SEM for the specimen annealed at 850 ◦ C for 24 h: (a) at the magnification of ×5000; (b) at the magnification of ×50,000.
FE-SEM micrographs of the fractures for the specimen annealed for 24 h. At low magnifications (×500, for instance) its fracture morphology (not shown here) is similar to that (Fig. 3a) for as-HIPed specimens, except for appearance of some porous like structure. However, its morphology at higher magnification (×5000) (Fig. 4a) is clearly different from that for as-HIPed specimen (Fig. 3b). One can see that the fracture surface is filled with dimples with their sizes being in the range of 0.2–2 m, as distinguished from the granular feature of the as-HIPed specimen (Fig. 3b). Fig. 4b gives the fractograph at much high magnification for the same fracture surface. This graph reveals that microscopically the annealed specimen displays very ductile fracture mode: mesoscale dimples appear with grains locating frequently inside. However, grain growth takes place substantially in this specimen, with some of them growing beyond 100 nm, although there are still some grains with sizes below 50 nm.
Tensile tests show that as-HIPed n-Ni–Fe fails prematurely in elastic region and its ultimate strength is only ∼1.1 GPa. This value is much smaller than the yield strength (∼1.6 GPa) obtained from compression tests [15] and that (∼1.5 GPa) evaluated from hardness measurements here, indicating asymmetry between tension and compression. This mechanical behavior is presumably related to the defects present in the specimens. Since our material is synthesized through powder fabrication and subsequent compaction by using HIPing, interfaces between agglomerate particles formed during HIPing should be weak places, where some microflaws or unclosed pores may exist due to residual unformed bonding. This speculation is reasonable, considering that the relative densities of the specimens are less than 100% of the theoretical one. On the other hand, deformation through dislocation activity in this material is expected to be difficult due to its small grain size (37 nm), considering that multiplication and slip of dislocations within these small grains need too much high stress. Furthermore, grain boundary sliding at room temperature can hardly happen in this material [15]. Therefore, under action of applied external stresses cracking would first initiate in the interfaces or existing microflaws there would propagate before any plastic deformation occurs, leading to premature failure of the material. The explanation of premature failure through interfacial cracking is consistent with the fractography, where the granular structures with the characteristic dimensions of agglomerate particle sizes appear (Fig. 3b). The observed vein-like patterns at low magnification (Fig. 3a) for as-HIPed specimen are seemingly similar to that observed in metallic glasses. But, their microscopic details are different. Vein patterns observed in metallic glass are usually featureless, and have no microstructural details [17]; while vein-like patterns observed here have, as mentioned above, nano-scale dimples. Since these vein-like patterns are presumably caused through tearing off, high rates of stain and cracking propagation along interfaces between particles are expected, which would give rise to local temperature rise. This would facilitate coalescence of free volumes, and so lead to formation of the nano-dimples. After the specimens are annealed at higher temperatures, fracture strength increases and hardness (yield stress) decreases, corresponding to which the fracture mode changes to more ductile manner. This change behavior should result from its microstructural and/or defect evolution during annealing. This change of microstructures may include: (1) grain growth; (2) enhancement of inter-particle adhesion; (3) void (porosity) formation coming from coalescence of residual micro-voids or free volumes. It is evident that grain growth (comparing Fig. 3c with Fig. 4b) may soften the material and enable the intra-granular deformation to become easier, leading to decrease in its yield stress (hardness) with annealing time (see Fig. 2). While inter-particle adhesion enhancement should inhibit inter-particle cracking
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and postpone occurrence of interfacial failure, which lead inevitably to increase in fracture strength with annealing, as observed in mechanical tests (Fig. 2). The large ((sub-) micrometer-sized) dimples observed (Fig. 4a) in the fractures of annealed specimens can be ascribed to the voids (porosity) formed through coalescence of residual micro-voids or free volumes. However, we note that ductile fracture morphology in the annealed specimens could be divided into two different levels: micrometer scale and mesoscale. Apart from the large dimples (Fig. 4a) being related to void formation, on much finer scale the formation of a large number of mesoscale dimples with sizes of several hundred nanometers (Fig. 4b) could be related to plastic deformation due to their configurations and dense population. They would initiate through stress concentrations produced at smaller nano-grains and develop through intragranular deformation in larger grains, for nano-grains are frequently observed to locate inside mesoscale-dimples (Fig. 4b). It is understandable that smaller grains with higher resistance to dislocation motion should act as obstacles to deformation. Plastic deformation would therefore proceed in large grains around the smaller grains. In other words, the mesoscale-dimples would presumably be associated with both concentrating effect of smaller grains and increasing deformation ability in larger grains. This speculation is supported by the fact that as annealing time reaches 24 h yield strength (the hardness Hv /3) is much smaller than fracture strength (Fig. 2) and macroscopic plastic deformation occurs substantially (Fig. 1). 5. Conclusions The characteristics of tensile deformation and fracture for as-prepared and annealed n-Ni–Fe alloy were investigated. The results indicate that the as-prepared specimen fails prematurely with ultimate strength (σ f ∼1.15 GPa) lower than yield strength evaluated from its Vickers microhardness. However, σ f increases and the hardness (yield strength) decreases monotonously with annealing time. Fractography indicates that for the as-prepared specimen inter-particle cracking is the main fracture mode, with nano-dimples appearing locally. When annealing time ta > ∼1.2 h σ f surpasses yield strength; correspondingly fractures are filled with a large number of both (sub-) micrometer-scale and mesoscale dimples as ta reaches 24 h. The improvement of fracture strength upon annealing can be ascribed mainly to
the enhancement of interfacial strength, while the decrease in yield strength (hardness) with annealing would originate from grain growth.
Acknowledgements The author (Qin) gratefully acknowledges Alexander von Humboldt Foundation for providing a fellowship. Specially he wishes to acknowledge Professor E. Nembach of Institut für Metallforschung, Universität Münster who contributed much to this work through arrangement of facilities and effective instructions. The authors are indebted to T. Krol for his assistance in mechanical tests and Bodycote IMT GmbH, Essen, Germany for HIPing experiments. Follow-up support from Academic Sinica through ‘Hundred Person Program’ and Korean Ministry of Science and Technology through the ‘2001 National Research Laboratory Program’ is gratefully acknowledged.
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