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Ultrafine composite microstructure in a bulk Ti alloy for high strength, strain hardening and tensile ductility
Acta Materialia 54 (2006) 1349–1357 www.actamat-journals.com
Ultrafine composite microstructure in a bulk Ti alloy for high strength, strain hardening and tensile ductility B.B. Sun a, M.L. Sui
a,*
, Y.M. Wang b, G. He c, J. Eckert d, E. Ma
b,*
a
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, Liaoning 110016, China b Department of Materials Science and Engineering, The Johns Hopkins University, Baltimore, MD 21218, USA c Department of Materials Science and Engineering, Shanghai Jiaotong University, Shanghai, China d Department of Materials and Geo-Sciences, Physical Metallurgy Division, Darmstadt University of Technology, Petersenstrasse 23, D-64287 Darmstadt, Germany Received 28 July 2005; received in revised form 1 November 2005; accepted 1 November 2005 Available online 4 January 2006
Abstract Ultrafine-grained (UFG) or nanostructured alloys usually lack the strain hardening capability needed to sustain uniform (tensile) deformation under high stresses. To circumvent this problem, we have designed a multi-phase composite microstructure in a Ti-based UFG alloy. The multi-component composition (Ti60Cu14Ni12Sn4Nb10) was chosen such that upon chill casting of the alloy the liquid underwent a metastable eutectic reaction, forming an in situ composite made of a micrometer-sized dendritic Ti-based solid solution intermixed with a UFG eutectic matrix. Such a microstructure imparts a high strength in excess of those of commercial Ti alloys, and, more importantly, allows strain hardening at relatively high rates. As a result, uniform elongation in tensile deformation was observed at high flow stresses. We present extensive microscopy results to illustrate the dislocation pile-ups and the origin of the high strength, as well as the extensive dislocation interactions and interface crossing responsible for the obvious strain hardening sustained to large plastic strains. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Ti alloy; Nanocomposite; Mechanical properties; Microstructure; Transmission electron microscopy
1. Introduction Nanostructures are often introduced into alloys for strengthening purposes. An obvious example is the extremely fine second-phase particles used for precipitation and dispersion hardening [1]. Recently, a large number of new nanostructured alloys and composites have been developed, in which the dominant (matrix) phase itself has a nanoscale or ultrafine-grained (UFG) structure [2]. While such materials offer high strength, they usually have to be made into bulk form either through multiple processing steps, e.g., via powder consolidation [3] and crystallization from amorphous precursors [4], or by employing *
Corresponding authors. E-mail addresses: [email protected] (M.L. Sui), [email protected] (E. Ma).
techniques that are not easily commercially viable, such as severe plastic deformation (SPD) using equal channel angular extrusion [5]. The powder route often suffers from porosity and contamination problems in the consolidation step [6], whereas the crystallization method imposes severe limitations on the sample size due to the limited dimensions of the amorphous precursors available. The crystallization products are also often intermetallic phases that can be rather brittle. With the SPD technique, due to the tremendous deformation needed to refine the microstructure, it is difficult to process very strong alloys. Moreover, upon deformation these high-strength nano/UFG materials rarely work harden at sufficiently high rates [6,7], making them prone to plastic instabilities. In uniaxial tensile deformation, the UFG alloys exhibit very small uniform tensile elongation. Therefore, it is of interest to develop a one-step
1359-6454/$30.00 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.11.011
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process that can directly lead to nano/UFG alloys in bulk form, e.g., through a casting operation, with an in situ formed microstructure that provides a good combination of desirable mechanical properties, in particular both high strength and tensile ductility. In this paper, we discuss the microstructure and mechanical properties of a bulk Ti60Cu14Ni12Sn4Nb10 (at.%) alloy based on UFG eutectic microstructure. In Section 2, we first present a general discussion of the microstructure design strategy. The experimental procedures are described in Section 3. Section 4 demonstrates the desirable mechanical properties of the alloy, including its high strength in combination with large plasticity, and in particular its strong strain hardening ability and tensile ductility. To correlate the properties with microstructure, a detailed structural and compositional analysis of the microstructures, before and after deformation, is given in Section 4. In the discussions presented in Section 5, the dislocation slip characteristics uncovered in microscopy studies are used to explain the strength, strain hardening, and ductility observed. The summary in Section 6 highlights the benefits that can be derived from tailoring such a UFG microstructure. 2. Microstructure design strategy To achieve a high strength, we desire a nano/ultrafinescale eutectic structure that can be obtained in liquid casting by properly choosing the alloy composition and controlling the solidification parameters. In a previous publication, we have presented an analysis of eutectic solidification to delineate the range of casting conditions that can make such UFG eutectic structures accessible in solidification of multi-component Ti-based liquid alloys [8]. It is important to note that together with the ultrafine eutectic microstructure, a ductile phase with a much larger grain size can be designed to be among the solidification product phases. An example is the dendrites of a primary solid solution grown before the eutectic reaction commences. One may then expect to reap the benefits of simultaneously coexisting high strength and large plasticity in a ‘‘bimodal’’ or composite microstructure, i.e., a combination of micrometer-sized dendrites and a UFG eutectic matrix [9,10]. To obtain such a solidification microstructure, we take advantage of a pseudo-binary system, Ti(Nb,Sn)–Cu(Ni). Elements such as Nb stabilize the high-temperature bodycentered cubic (bcc) Ti solution phase to low temperatures. A metastable eutectic reaction (L solidifies into b-Ti solution and c-CuTi) occurs in solidification, as the formation of the equilibrium CuTi2 intermetallic would be delayed due to its complicated crystal structure, which is difficult to form as a result of the confusion and the slow kinetics brought about by the multi-component recipe, suppressing its growth [8]. We selected a composition of Ti60Cu14Ni12Sn4Nb10. As is known from a calculated phase diagram showing the metastable equilibrium between the bcc Ti solution and the c-TiCu phase [8], there will be a rather deep metastable eutectic with a plunging liquidus line. To
obtain the bcc Ti as the primary phase, our composition will be on the bcc Ti side of the eutectic composition, and as such the solidification products would produce the desired microstructure as discussed in the preceding paragraph. The multiple elements added also render the eutectic growth more sluggish, favoring the refinement of lamellar spacing to ultrafine scales. For more discussions of solidification in such systems, the reader is referred to Ref. [8]. 3. Experimental Ingots with the desired composition (Ti60Cu14Ni12Sn4Nb10) were prepared by arc melting a mixture of pure elements in a water-cooled copper hearth in an argon atmosphere. The ingots were then re-melted in a quartz crucible and injected into a copper mould to cast cylinders 2 mm in diameter. The microstructures of this alloy were characterized by means of X-ray diffraction (XRD) analysis, scanning electron microscopy (SEM), transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM). XRD measurements were carried out using a Rigaku D/max 2400 X-ray diffractometer with Cu Ka radiation (k = 0.15405 nm). A Philips XL30FEG SEM equipped with an energy-dispersive spectrometer (EDS) was used for microstructure observations and composition analyses. The samples for SEM observations were etched in a solution of HF and HNO3 at room temperature. JEM 2000 FX and JEM 2010 instruments were used for TEM and HRTEM observations. TEM samples were prepared using standard procedures involving ion milling. For mechanical property assessments, both compressive tests and tensile tests were carried out at a strain rate of 1 · 104/s at room temperature. 4. Results 4.1. Mechanical properties We present the attractive mechanical properties first to aid in the detailed understanding of the microstructure and dislocation behavior to be discussed in later sections. The typical compressive and tensile true stress–strain curves are displayed in Fig. 1. The 0.2% offset yield strength is 1.2 GPa, higher than that of commercial Ti alloys. The alloy shows pronounced work hardening in both compression and tensile tests. As shown in the inset of Fig. 1, the work hardening rate is exceptionally large at the high stress levels. An impressive ultimate strength of 1.75 GPa is reached in compression. The strong working hardening allows uniform elongation in tension, leading to a fast ascending curve in the tensile test, unlike most UFG and nanocrystalline metals and alloys that show plunging tensile curves peaking very early in plastic deformation [11]. The plastic strain to failure reached 10.3% in compression. The obvious tensile ductility is rarely observed in multi-component alloys that form metallic glasses or nanocrystalline phases.
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Fig. 2. XRD trace of the Ti60Cu14Ni12Sn4Nb10 alloy. Fig. 1. Room-temperature compressive and tensile true stress–strain curves of the Ti60Cu14Ni12Sn4Nb10 alloy. The work hardening rate, where r is the true stress and e the true strain, is defined here as H ¼ r1 or oe e_ plotted in the inset.
The tensile elongation terminated after 3.2% of plastic strain; but this is not because of necking. We noticed that, as shown in the inset of Fig. 1, at the point where tensile failure started, the strain hardening rate is still a factor of 4 away from the Conside`re instability criterion for the onset of necking (H 6 1, see caption of Fig. 1). The presence of an appreciable strain hardening rate suggests that the tensile failure is not preceded by strain localization when the plastic instability point is reached. Rather the failure is believed to be triggered by flaw(s) in the sample that served as crack nucleation sites under the high stresses. The most likely source of flaws is the casting pores, which were indeed observed by inspecting the cast ingot. The important message from Fig. 1 (especially the compressive behavior) is that the alloy is intrinsically very ductile, and the strain hardening capacity is rather large so that deformation can be stabilized to large strains without the onset of localized deformation concentrating large strains. In fact, its working hardening rate exceeds by far those of recently developed alloys of similar strength, such as nanocrystalline and amorphous metals and alloys [12–15]. Eliminating solidification flaws would be important for achieving large tensile ductility, but this task is beyond the scope of the present study. We predict large tensile ductility after the casting processing is improved in future studies.
other peaks are present. A careful matching with known standards reveals that these peaks are from c-CuTi and CuTi2 intermetallic phases, as indicated in Fig. 2. We show later that the CuTi phase is involved in the metastable eutectic reaction, whereas some CuTi2 forms in the late stages of solidification. The XRD peaks are significantly broadened, indicating small grain/domain sizes. Also, trace amounts of TiC due to contaminations are present in this alloy. The TiC particles in fact often serve as heterogeneous sites for the nucleation of the b-Ti solution primary phase. Fig. 3 shows an SEM image of the Ti60Cu14Ni12Sn4Nb10 alloy taken near the center of the cross-section of the ingot. The micrograph shows that the alloy is a composite of micrometer-sized dendrites dispersed in a matrix. The rounded dendrite is a few micrometers in length. The arm spacing of the dendrite is less than 1 lm. Our EDS results show that the dendritic phase is enriched in Ti, Nb and Sn.
4.2. Microstructures 4.2.1. Composite features Fig. 2 shows the XRD trace of the alloy under investigation. The predominant phase identified is a bcc b-Ti solid solution. The lattice parameter of this solid solution phase is estimated to be 0.327 nm from the strongest peak. In addition to the diffraction peaks of b-Ti solid solution,
Fig. 3. SEM image of the Ti60Cu14Ni12Sn4Nb10 alloy, taken near the cross-section center of the ingot, showing the dendrite and matrix morphology.
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This dendritic phase has bcc structure according to the phase identification by XRD and the selected area electron diffraction (SAED) analysis by TEM. It is thus obvious that the dendritic phase is a b-Ti(Nb,Sn) solid solution. It
Fig. 4. Bright-field TEM image of the Ti60Cu14Ni12Sn4Nb10 alloy: (a) dendrites and eutectic colonies with ultrafine structure; (b) the region between eutectic colonies. The inset shows the [3 3 1] diffraction of CuTi2 phase.
is in fact well known that Nb is a strong b stabilizer in Ti alloys. The matrix is dominated by micrometer-sized colonies with ultrafine lamellar structure inside, as the product of the eutectic reaction. From EDS analysis, the colonies are enriched in Ti, Cu and Ni. As such, they correspond to the c-CuTi and/or CuTi2 phase in the XRD trace. The growth boundaries of the colonies are clearly observed. The bright-field TEM image in Fig. 4(a) shows the dendrites, and the eutectic colonies with ultrafine structure. Fig. 4(b) shows the region in between eutectic colonies, where the CuTi2 phase is located. The presence of CuTi2 was confirmed by SAED (see inset in Fig. 4(b)). It should be noted that the CuTi2 phase was only found in the regions between the colonies, rather than inside the colonies. To identify the phases inside the eutectic colony, SAED was applied. In the SAED pattern in Fig. 5(a), one set of the spots is indexed as the diffraction pattern of the [1 1 1] zone axis of the b-Ti solid solution, and the other set of the spots is identified as the diffraction pattern of the [2 2 1] zone axis of the c-CuTi intermetallic. In Fig. 5(b), the two sets of electron diffraction spots are identified as [2 0 1] of b-Ti and [4 0 1] of c-CuTi, respectively. The corresponding orientation relationship between the b and c phases is [1 1 1]b// [2 2 1]c, ð1 1 0Þb ==ð1 1 0Þc . The SAED results confirm the existence of the c-CuTi phase as suggested by XRD and reveal the coexistence of the two phases, b-Ti and c-CuTi, in the eutectic colony. The TEM observations further demonstrate that inside the colony ultrafine b/c phases are arranged in an alternating fashion due to eutectic solidification in the Ti(Nb,Sn)–Cu(Ni,Sn) pseudo-binary system. Also, a graded microstructure along the radial direction was observed in this alloy, as shown in Fig. 6. Near the surface of the ingot (Fig. 6(a) and (b)), the matrix microstructure cannot be resolved using SEM and the size of the dendrite is small. By use of TEM imaging (Fig. 6(b)), the average size scale of the dendrites, eutectic colony and lamellar spacing is determined to be about 0.5 lm, 2 lm
Fig. 5. SAED patterns of the colony. Along (a) the [1 1 1] zone axis and (b) the [2 0 1] zone axis of the b-Ti phase.
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Fig. 6. SEM and TEM images showing the size gradient along the radial direction of the ingot. (a) SEM image of a structure about 100 lm from the surface of the ingot and (b) its corresponding TEM image. (c) SEM image of a structure close to center of the ingot and (d) its corresponding TEM image.
and 130 nm, respectively. In contrast, close to the center of the ingot (see Fig. 6(c) and (d)), these three length scales increase to 1–1.5 lm, 4 lm and 100–300 nm, respectively. This is an interesting phenomenon, as one can achieve a grain-size-graded microstructure. The observation that this happens even across a small cross-section suggests that the structure of this alloy is very sensitive to the cooling rate. Based on the results above, the as-cast Ti60Cu14Ni12Sn4Nb10 alloy rod can be described as a composite, with micrometer-sized b-Ti primary solid solution dendrites dispersed in a UFG eutectic b/c matrix, with a length scale (grain size) gradient along the radial direction. 4.2.2. Structure of the ultrafine eutectic matrix Fig. 7(a) shows a TEM image of the ultrafine eutectic microstructure. The first impression is that the b and c phases exhibit a fine lamellar eutectic structure. However, when observed along another crystallographic direction, Fig. 7(b), the cross-section of the b phase is embedded in the c phase and has the shape of a particle. These observations indicate that the eutectic is actually rod-like (b phase). A statistical analysis shows that on average the dimensions of the b phase inside the eutectic matrix are about a few micrometers in length and 70–80 nm in the transverse direction. The volume fraction of the b phase in the eutectic matrix is estimated to be about 23% based on image analysis of the area fractions observed. To summarize, the
Fig. 7. Bright-field TEM images of the eutectic matrix observed along near (a) [2 0 1]b and (b) [1 1 1]b.
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eutectic matrix in the Ti60Cu14Ni12Sn4Nb10 alloy has a rod eutectic structure on an ultrafine scale, i.e., the b rods are embedded in the c matrix. The b/c interfaces have also been investigated by means of HRTEM observations. The two phases form a semicoherent interface. Fig. 8(a) and (b) shows a HRTEM image and the corresponding Fourier filtered image of the b/c interface taken along the [1 1 1]b direction. Some misfit dislocations, along with dislocations of other types, are observed at the well-bonded interface. Fig. 9(a) and (b) are the original and the Fourier filtered HRTEM images of the b phase inside the nano/ultrafine eutectic matrix along the [1 1 1]b direction. The inset is a Fourier transformed pattern of the HRTEM image, in which obviously elongated spots can be observed. In Fig. 9(b), stacks corresponding to the elongated spots in the inset are
Fig. 9. (a) Original and (b) the Fourier filtered HRTEM images of the b phase inside the nano/ultrafine eutectic matrix along the [1 1 1]b direction. The inset shows a Fourier transformed pattern of the HRTEM image.
observed; also, dislocations and lattice distortions are found in the b phase. These observations suggest that many lattice defects are stored in the as-cast eutectic matrix. The microstructural features uncovered in this section are correlated with mechanical properties in the following section.
Fig. 8. (a) HRTEM image and (b) the corresponding Fourier filtered image of the b-Ti/c-CuTi interface in the eutectic matrix.
4.2.3. Structure after deformation Fig. 10 shows the SEM image for the alloy after the compression test. During the uniaxial compression the load was applied in the horizontal direction. Some deformation bands (denoted by black arrows) are clearly seen, which are often described as shear bands [16]. The spacing of these bands is about 1–2 lm. These shear bands travel across rather long distances, propagating through dendritic arms and the eutectic colonies, roughly along the direction 45
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Fig. 10. SEM secondary electron image of the surface of the deformed alloy, showing the cross-colony shear bands (black arrows) and the intracolony slip bands (white arrows). The inset shows an SEM backscattered electron image of the region enclosed by the dashed rectangle, showing the dendrite arms and the propagation of the slip bands.
to the compression loading direction (the maximum shear stress direction). Besides these shear bands, there are shorter bands that appear to be of a high density (denoted by white arrows). Most of these bands propagate inside the
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dendrites and across the gap (matrix) in between the closely spaced dendritic arms, see the inset in Fig. 10. The spacing of these bands is less than 700 nm, and the direction of these bands has no special relationship with the compression direction, indicating that these are slip bands. By use of TEM, the major shear bands that travel long distances such as those in Fig. 10 were not observed. They are neither very large in number nor in width, making them difficult to be caught in TEM sample preparation. However, the slip bands, which we believe contribute the most to plastic deformation, were easily found. From Fig. 11(a), the spacing of the slip bands is estimated to be in the range 90–300 nm. The value is in reasonable agreement with that obtained from Fig. 10. Two slip systems were active. By a series of tilting experiments in the TEM instrument, the two systems were determined to be ð1 1 2Þ½1 1 1 and ð1 1 2Þ½1 1 1, as shown in Fig. 11(a). In a given ‘‘dendrite colony’’ (a dendrite with arms), the slip bands extend along virtually the same direction. The extensive slip deformation induced pile-ups of dislocations at the interfaces between the dendrite arm and the matrix, see Fig. 11(b) and (c). The dislocation pile-ups, as often invoked in the interpretation of the Hall–Petch relationship, lead to stress concentration and thus the transmission of the slip across the interface into the matrix, see Fig. 11(d). The steps left at the interface due to such crossing events are shown in Fig. 11(b), denoted by black arrows. The matrix between
Fig. 11. Bright-field TEM images of the deformed alloy. (a) Slip bands in each ‘‘dendrite colony’’ aligned approximately along the same direction; (b) pileup of dislocations and steps (black arrows) left by slip across the interface; (c) enlarged view of the dislocation pile-up; (d) slip band propagating into the harder matrix phase.
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the dendrite arms is mostly the c-CuTi phase, again in semicoherent relationship with the b dendrite arms. As such, the slip band runs straight across the interface, Fig. 11(d). 5. Discussion The compressive and tensile stress–strain curves of the Ti60Cu14Ni12Sn4Nb10 alloy shown in Fig. 1 demonstrate several desirable properties that can be linked to its unique microstructure. The mechanical properties of composites are sometimes interpreted using the rule of mixtures [17] P c ¼ P dV d þ P mV m
ð1Þ
where Pc is the property of the composite, and Pd and Pm are the properties of the dendrite and matrix, respectively. Vd andVm are the volume fractions of the dendrite and matrix, respectively. In our alloy, the dendrite b phase is the soft phase while the nano/ultrafine eutectic matrix is the hard phase. The volume fraction of dendrite is estimated to be approximately 50% based on image analysis. As expected from Eq. (1), the yield strength shown in Fig. 1 would be lower than that in Ref. [8], in which the volume fraction of the softer dendrite is about 20% and the yield strength of that composite is correspondingly higher (at 1.5 GPa). Obviously, the property desired can be adjusted by simply tailoring the volume fractions of the participating phases, when designing the composition of the alloy [8,18]. Even with the high volume faction of b-Ti, our alloy still exhibits a yield strength higher than commercial Ti alloys [19]. The high strength of the hard matrix comes from several microstructural features discussed in Section 4.2, including the UFG eutectic structure, the high contents of multiple solutes that lead to solution hardening, and the lattice defects in the as-cast alloy (i.e., high densities of dislocations, faults and lattice distortions revealed in TEM as described above (Fig. 9)). That the UFG eutectic is the strengthening phase can be easily appreciated by the observation that dislocations pile-up at the dendrite/matrix boundary before they enter the matrix from the dendrites. The trace amount of TiC phase, an unexpected byproduct due to contamination during sample preparation, possibly contributes as another strengthening phase. The results of Section 4.2.3 suggest that the numerous slip bands due to dislocation activities in the much larger ductile dendrites contribute to the overall ductility of this alloy. Furthermore, the accumulation of dislocations and interactions of the different slip systems contribute to the obvious strain hardening. With escalating stresses, the slip band transmits across the interfaces and propagates in the harder matrix. Without the ductile reinforcements, monolithic UFG alloys usually cannot sustain high rates of strain hardening at high stress levels, making it difficult to attain uniform tensile ductility [20]. As mentioned before, in addition to the slip bands that propagate within a dendrite colony, there are also shear bands that travel past many dendrite colonies as well as eutectic colonies. These sporadic shear bands, observed in
SEM, would also contribute to plastic strain to some extent, but not to the degree of that discussed in Section 4.2.3 for the plasticity inside the dendrite colonies. The eutectic structure, especially when decorated with a high density of defects as depicted in Fig. 9, is expected to be very strong when compared with the dendrite phase. Their contribution to the strength is expected, but their contribution to plastic strain and strain hardening is not expected to be very large, just as in the case for nano/UFG multilayers [21]. The graded length scales of the features (grain size, arm spacing, etc.) along the radial direction observed in this cast alloy can play a positive role in providing useful properties. A similar grain-size-graded structure has been achieved by means of surface mechanical attrition treatment [22], and was found to impart beneficial properties. Hanlon et al. [23] recently studied the fatigue response of nanocrystalline and UFG metals and alloys. They found that grain refinement generally leads to an increase in the resistance to failure under stress-controlled fatigue. A high-strength nanostructure is desirable at the surface where the cracks initiate. However, Hanlon et al. also concluded that nanostructures would have a generally deleterious effect on the resistance to the growth of a fatigue crack once it starts. In other words, one would prefer less refined structure in the interior in order to stop crack growth. In this context, the micro/nanostructure gradient we observed naturally offers a functionally graded material for such applications. 6. Summary A Ti60Cu14Ni12Sn4Nb10 alloy has been processed to be composed mostly of a micrometer-sized dendritic Ti-based bcc solid solution uniformly dispersed within a UFG matrix. The matrix is characterized by a rod eutectic microstructure. As the system is approximately a pseudo-binary undergoing a metastable eutectic reaction, the eutectic phases are primarily a b-Ti(Nb,Sn) solid solution alternating with a c-TiCu(Ni) intermetallic compound, with the b/c orientation relationship being [1 1 1]b//[2 2 1]c and ð110Þb ==ð110Þc . The dendrites provide the large plastic strains and strain hardening (e.g., dislocation storage) capability, whereas the strength mainly comes from the matrix where the UFG eutectic structure and the high concentrations of multiple solutes and crystallographic defects strengthen the alloy. We obtained from TEM clear and direct evidence of the microscopic dislocation mechanisms, in terms of dislocation pile-ups at the strengthening interfaces before the dislocations come from the softer phase and enter the harder phase, as well as extensive build-up of networks of slip bands in the dendrite arms. The transmission of slip across the interfaces (both dendrite/matrix and the semi-coherent boundaries between the eutectic layers) requires high stresses, hence leading to high strength. The slip propagation also produces slip steps and hence large plastic strains, and promotes work hardening due to the interactions among the accumulating dislocations and
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multiple activated slip systems. The numerous interphase boundaries in the microstructure accentuate the role of interfaces in the strengthening and ductilization of the alloy. This work demonstrates one strategy for an alloy to achieve high strength, large (tensile) ductility, high strain hardening rates and functionally graded properties, all at the same time. While it is also possible to obtain such properties in an elemental [9] or simple binary system [24], we have shown that a multi-component recipe offers more ‘‘knobs to turn’’. In this particular Ti alloy case, the addition of Nb stabilizes the b phase, allowing the ductile bcc b to be the primary phase (rather than the hexagonal close packed a). It also effectively creates a relatively deep L ! b + c metastable eutectic, promoting the formation of a UFG eutectic structure. The multiple solutes also suppress the diffusion and coarsening of the eutectic structure, and are also helpful for strengthening purposes. Our in situ composite route can also yield useful materials in bulk form via a one-step casting process.
Acknowledgments The authors acknowledge the support of the National Natural Science Foundation of China (Grant Nos. 50125103 and 50323009), the Shenyang Center for Interfacial Materials (the MANS research team), the Alexander von Humboldt Foundation, and the EU within the framework of the RTN-Network on bulk metallic glasses (HPRN-CT-2000-00033).
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References [1] Dieter GE. In: Mechanical metallurgy. Boston: McGraw-Hill; 1986. p. 212. [2] Mccandlish LE, Kear BH, Kim BK. Mater Sci Technol 1990;6:953. [3] He L, Ma E. Nanostructured Mater 1996;7:327; He L, Ma E. J Mater Res 1996;11:72. [4] Lu K, Wang JT, Wei WD. J Appl Phys 1991;69:522. [5] Jia D, Wang YM, Ramesh KT, Ma E, Zhu YT, Valiev RZ. Appl Phys Lett 2001;79:611. [6] Koch CC. Scripta Mater 2003;49:657. [7] Tellkamp VL, Melmed A, Lavernia EJ. Metall Mater Trans A 2001;32:2335. [8] Dai QL, Sun BB, Li Y, He G, Eckert J, Ma E, et al. J Mater Res 2004;19:2557. [9] Wang YM, Chen M, Zhou F, Ma E. Nature 2002;419:912. [10] He G, Eckert J, Lo¨ser W, Schultz L. Nature Mater 2003;2:33. [11] Ma E. Scripta Mater 2003;49:663. [12] Wang YM, Ma E. Acta Mater 2004;52:1699. [13] Guo FQ, Wang HJ, Poon SJ, Shiflet GJ. Appl Phys Lett 2005;86:091907. [14] Youssef KM, Scattergood RO, Murty KL, Horton JA, Koch CC. Appl Phys Lett 2005;87:091904. [15] Cheng S, Ma E, Wang YM, Kecskes LJ, Youssef KM, Koch CC, et al. Acta Mater 2005;53:1521. [16] Kim KB, Das J, Baier F, Eckert J. Appl Phys Lett 2005;86:171909. [17] Chadwick GA. Metal Sci J 1975;9:300. [18] He G, Eckert J, Lo¨ser W, Hagiwara M. Acta Mater 2004;52:3035. [19] Manero JM, Gil FJ, Planell JA. Acta Mater 2000;48:3353. [20] Ma E. Nature Mater 2003;2:7. [21] Misra A, Hirth JP, Hoagland RG, Embury JD, Kung H. Acta Mater 2004;52:2387. [22] Tao NR, Zhang HW, Lu J, Lu K. Mater Trans 2003;44:1919. [23] Hanlon T, Kwon YN, Suresh S. Scripta Mater 2003;49:675. [24] Louzguine DV, Kato H, Louzguina LV, Inoue A. J Mater Res 2004;19:3600.
Ultrafine composite microstructure in a bulk Ti alloy for high strength, strain hardening and tensile ductility B.B. Sun a, M.L. Sui
a,*
, Y.M. Wang b, G. He c, J. Eckert d, E. Ma
b,*
a
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, Liaoning 110016, China b Department of Materials Science and Engineering, The Johns Hopkins University, Baltimore, MD 21218, USA c Department of Materials Science and Engineering, Shanghai Jiaotong University, Shanghai, China d Department of Materials and Geo-Sciences, Physical Metallurgy Division, Darmstadt University of Technology, Petersenstrasse 23, D-64287 Darmstadt, Germany Received 28 July 2005; received in revised form 1 November 2005; accepted 1 November 2005 Available online 4 January 2006
Abstract Ultrafine-grained (UFG) or nanostructured alloys usually lack the strain hardening capability needed to sustain uniform (tensile) deformation under high stresses. To circumvent this problem, we have designed a multi-phase composite microstructure in a Ti-based UFG alloy. The multi-component composition (Ti60Cu14Ni12Sn4Nb10) was chosen such that upon chill casting of the alloy the liquid underwent a metastable eutectic reaction, forming an in situ composite made of a micrometer-sized dendritic Ti-based solid solution intermixed with a UFG eutectic matrix. Such a microstructure imparts a high strength in excess of those of commercial Ti alloys, and, more importantly, allows strain hardening at relatively high rates. As a result, uniform elongation in tensile deformation was observed at high flow stresses. We present extensive microscopy results to illustrate the dislocation pile-ups and the origin of the high strength, as well as the extensive dislocation interactions and interface crossing responsible for the obvious strain hardening sustained to large plastic strains. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Ti alloy; Nanocomposite; Mechanical properties; Microstructure; Transmission electron microscopy
1. Introduction Nanostructures are often introduced into alloys for strengthening purposes. An obvious example is the extremely fine second-phase particles used for precipitation and dispersion hardening [1]. Recently, a large number of new nanostructured alloys and composites have been developed, in which the dominant (matrix) phase itself has a nanoscale or ultrafine-grained (UFG) structure [2]. While such materials offer high strength, they usually have to be made into bulk form either through multiple processing steps, e.g., via powder consolidation [3] and crystallization from amorphous precursors [4], or by employing *
Corresponding authors. E-mail addresses: [email protected] (M.L. Sui), [email protected] (E. Ma).
techniques that are not easily commercially viable, such as severe plastic deformation (SPD) using equal channel angular extrusion [5]. The powder route often suffers from porosity and contamination problems in the consolidation step [6], whereas the crystallization method imposes severe limitations on the sample size due to the limited dimensions of the amorphous precursors available. The crystallization products are also often intermetallic phases that can be rather brittle. With the SPD technique, due to the tremendous deformation needed to refine the microstructure, it is difficult to process very strong alloys. Moreover, upon deformation these high-strength nano/UFG materials rarely work harden at sufficiently high rates [6,7], making them prone to plastic instabilities. In uniaxial tensile deformation, the UFG alloys exhibit very small uniform tensile elongation. Therefore, it is of interest to develop a one-step
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process that can directly lead to nano/UFG alloys in bulk form, e.g., through a casting operation, with an in situ formed microstructure that provides a good combination of desirable mechanical properties, in particular both high strength and tensile ductility. In this paper, we discuss the microstructure and mechanical properties of a bulk Ti60Cu14Ni12Sn4Nb10 (at.%) alloy based on UFG eutectic microstructure. In Section 2, we first present a general discussion of the microstructure design strategy. The experimental procedures are described in Section 3. Section 4 demonstrates the desirable mechanical properties of the alloy, including its high strength in combination with large plasticity, and in particular its strong strain hardening ability and tensile ductility. To correlate the properties with microstructure, a detailed structural and compositional analysis of the microstructures, before and after deformation, is given in Section 4. In the discussions presented in Section 5, the dislocation slip characteristics uncovered in microscopy studies are used to explain the strength, strain hardening, and ductility observed. The summary in Section 6 highlights the benefits that can be derived from tailoring such a UFG microstructure. 2. Microstructure design strategy To achieve a high strength, we desire a nano/ultrafinescale eutectic structure that can be obtained in liquid casting by properly choosing the alloy composition and controlling the solidification parameters. In a previous publication, we have presented an analysis of eutectic solidification to delineate the range of casting conditions that can make such UFG eutectic structures accessible in solidification of multi-component Ti-based liquid alloys [8]. It is important to note that together with the ultrafine eutectic microstructure, a ductile phase with a much larger grain size can be designed to be among the solidification product phases. An example is the dendrites of a primary solid solution grown before the eutectic reaction commences. One may then expect to reap the benefits of simultaneously coexisting high strength and large plasticity in a ‘‘bimodal’’ or composite microstructure, i.e., a combination of micrometer-sized dendrites and a UFG eutectic matrix [9,10]. To obtain such a solidification microstructure, we take advantage of a pseudo-binary system, Ti(Nb,Sn)–Cu(Ni). Elements such as Nb stabilize the high-temperature bodycentered cubic (bcc) Ti solution phase to low temperatures. A metastable eutectic reaction (L solidifies into b-Ti solution and c-CuTi) occurs in solidification, as the formation of the equilibrium CuTi2 intermetallic would be delayed due to its complicated crystal structure, which is difficult to form as a result of the confusion and the slow kinetics brought about by the multi-component recipe, suppressing its growth [8]. We selected a composition of Ti60Cu14Ni12Sn4Nb10. As is known from a calculated phase diagram showing the metastable equilibrium between the bcc Ti solution and the c-TiCu phase [8], there will be a rather deep metastable eutectic with a plunging liquidus line. To
obtain the bcc Ti as the primary phase, our composition will be on the bcc Ti side of the eutectic composition, and as such the solidification products would produce the desired microstructure as discussed in the preceding paragraph. The multiple elements added also render the eutectic growth more sluggish, favoring the refinement of lamellar spacing to ultrafine scales. For more discussions of solidification in such systems, the reader is referred to Ref. [8]. 3. Experimental Ingots with the desired composition (Ti60Cu14Ni12Sn4Nb10) were prepared by arc melting a mixture of pure elements in a water-cooled copper hearth in an argon atmosphere. The ingots were then re-melted in a quartz crucible and injected into a copper mould to cast cylinders 2 mm in diameter. The microstructures of this alloy were characterized by means of X-ray diffraction (XRD) analysis, scanning electron microscopy (SEM), transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM). XRD measurements were carried out using a Rigaku D/max 2400 X-ray diffractometer with Cu Ka radiation (k = 0.15405 nm). A Philips XL30FEG SEM equipped with an energy-dispersive spectrometer (EDS) was used for microstructure observations and composition analyses. The samples for SEM observations were etched in a solution of HF and HNO3 at room temperature. JEM 2000 FX and JEM 2010 instruments were used for TEM and HRTEM observations. TEM samples were prepared using standard procedures involving ion milling. For mechanical property assessments, both compressive tests and tensile tests were carried out at a strain rate of 1 · 104/s at room temperature. 4. Results 4.1. Mechanical properties We present the attractive mechanical properties first to aid in the detailed understanding of the microstructure and dislocation behavior to be discussed in later sections. The typical compressive and tensile true stress–strain curves are displayed in Fig. 1. The 0.2% offset yield strength is 1.2 GPa, higher than that of commercial Ti alloys. The alloy shows pronounced work hardening in both compression and tensile tests. As shown in the inset of Fig. 1, the work hardening rate is exceptionally large at the high stress levels. An impressive ultimate strength of 1.75 GPa is reached in compression. The strong working hardening allows uniform elongation in tension, leading to a fast ascending curve in the tensile test, unlike most UFG and nanocrystalline metals and alloys that show plunging tensile curves peaking very early in plastic deformation [11]. The plastic strain to failure reached 10.3% in compression. The obvious tensile ductility is rarely observed in multi-component alloys that form metallic glasses or nanocrystalline phases.
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Fig. 2. XRD trace of the Ti60Cu14Ni12Sn4Nb10 alloy. Fig. 1. Room-temperature compressive and tensile true stress–strain curves of the Ti60Cu14Ni12Sn4Nb10 alloy. The work hardening rate, where r is the true stress and e the true strain, is defined here as H ¼ r1 or oe e_ plotted in the inset.
The tensile elongation terminated after 3.2% of plastic strain; but this is not because of necking. We noticed that, as shown in the inset of Fig. 1, at the point where tensile failure started, the strain hardening rate is still a factor of 4 away from the Conside`re instability criterion for the onset of necking (H 6 1, see caption of Fig. 1). The presence of an appreciable strain hardening rate suggests that the tensile failure is not preceded by strain localization when the plastic instability point is reached. Rather the failure is believed to be triggered by flaw(s) in the sample that served as crack nucleation sites under the high stresses. The most likely source of flaws is the casting pores, which were indeed observed by inspecting the cast ingot. The important message from Fig. 1 (especially the compressive behavior) is that the alloy is intrinsically very ductile, and the strain hardening capacity is rather large so that deformation can be stabilized to large strains without the onset of localized deformation concentrating large strains. In fact, its working hardening rate exceeds by far those of recently developed alloys of similar strength, such as nanocrystalline and amorphous metals and alloys [12–15]. Eliminating solidification flaws would be important for achieving large tensile ductility, but this task is beyond the scope of the present study. We predict large tensile ductility after the casting processing is improved in future studies.
other peaks are present. A careful matching with known standards reveals that these peaks are from c-CuTi and CuTi2 intermetallic phases, as indicated in Fig. 2. We show later that the CuTi phase is involved in the metastable eutectic reaction, whereas some CuTi2 forms in the late stages of solidification. The XRD peaks are significantly broadened, indicating small grain/domain sizes. Also, trace amounts of TiC due to contaminations are present in this alloy. The TiC particles in fact often serve as heterogeneous sites for the nucleation of the b-Ti solution primary phase. Fig. 3 shows an SEM image of the Ti60Cu14Ni12Sn4Nb10 alloy taken near the center of the cross-section of the ingot. The micrograph shows that the alloy is a composite of micrometer-sized dendrites dispersed in a matrix. The rounded dendrite is a few micrometers in length. The arm spacing of the dendrite is less than 1 lm. Our EDS results show that the dendritic phase is enriched in Ti, Nb and Sn.
4.2. Microstructures 4.2.1. Composite features Fig. 2 shows the XRD trace of the alloy under investigation. The predominant phase identified is a bcc b-Ti solid solution. The lattice parameter of this solid solution phase is estimated to be 0.327 nm from the strongest peak. In addition to the diffraction peaks of b-Ti solid solution,
Fig. 3. SEM image of the Ti60Cu14Ni12Sn4Nb10 alloy, taken near the cross-section center of the ingot, showing the dendrite and matrix morphology.
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This dendritic phase has bcc structure according to the phase identification by XRD and the selected area electron diffraction (SAED) analysis by TEM. It is thus obvious that the dendritic phase is a b-Ti(Nb,Sn) solid solution. It
Fig. 4. Bright-field TEM image of the Ti60Cu14Ni12Sn4Nb10 alloy: (a) dendrites and eutectic colonies with ultrafine structure; (b) the region between eutectic colonies. The inset shows the [3 3 1] diffraction of CuTi2 phase.
is in fact well known that Nb is a strong b stabilizer in Ti alloys. The matrix is dominated by micrometer-sized colonies with ultrafine lamellar structure inside, as the product of the eutectic reaction. From EDS analysis, the colonies are enriched in Ti, Cu and Ni. As such, they correspond to the c-CuTi and/or CuTi2 phase in the XRD trace. The growth boundaries of the colonies are clearly observed. The bright-field TEM image in Fig. 4(a) shows the dendrites, and the eutectic colonies with ultrafine structure. Fig. 4(b) shows the region in between eutectic colonies, where the CuTi2 phase is located. The presence of CuTi2 was confirmed by SAED (see inset in Fig. 4(b)). It should be noted that the CuTi2 phase was only found in the regions between the colonies, rather than inside the colonies. To identify the phases inside the eutectic colony, SAED was applied. In the SAED pattern in Fig. 5(a), one set of the spots is indexed as the diffraction pattern of the [1 1 1] zone axis of the b-Ti solid solution, and the other set of the spots is identified as the diffraction pattern of the [2 2 1] zone axis of the c-CuTi intermetallic. In Fig. 5(b), the two sets of electron diffraction spots are identified as [2 0 1] of b-Ti and [4 0 1] of c-CuTi, respectively. The corresponding orientation relationship between the b and c phases is [1 1 1]b// [2 2 1]c, ð1 1 0Þb ==ð1 1 0Þc . The SAED results confirm the existence of the c-CuTi phase as suggested by XRD and reveal the coexistence of the two phases, b-Ti and c-CuTi, in the eutectic colony. The TEM observations further demonstrate that inside the colony ultrafine b/c phases are arranged in an alternating fashion due to eutectic solidification in the Ti(Nb,Sn)–Cu(Ni,Sn) pseudo-binary system. Also, a graded microstructure along the radial direction was observed in this alloy, as shown in Fig. 6. Near the surface of the ingot (Fig. 6(a) and (b)), the matrix microstructure cannot be resolved using SEM and the size of the dendrite is small. By use of TEM imaging (Fig. 6(b)), the average size scale of the dendrites, eutectic colony and lamellar spacing is determined to be about 0.5 lm, 2 lm
Fig. 5. SAED patterns of the colony. Along (a) the [1 1 1] zone axis and (b) the [2 0 1] zone axis of the b-Ti phase.
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Fig. 6. SEM and TEM images showing the size gradient along the radial direction of the ingot. (a) SEM image of a structure about 100 lm from the surface of the ingot and (b) its corresponding TEM image. (c) SEM image of a structure close to center of the ingot and (d) its corresponding TEM image.
and 130 nm, respectively. In contrast, close to the center of the ingot (see Fig. 6(c) and (d)), these three length scales increase to 1–1.5 lm, 4 lm and 100–300 nm, respectively. This is an interesting phenomenon, as one can achieve a grain-size-graded microstructure. The observation that this happens even across a small cross-section suggests that the structure of this alloy is very sensitive to the cooling rate. Based on the results above, the as-cast Ti60Cu14Ni12Sn4Nb10 alloy rod can be described as a composite, with micrometer-sized b-Ti primary solid solution dendrites dispersed in a UFG eutectic b/c matrix, with a length scale (grain size) gradient along the radial direction. 4.2.2. Structure of the ultrafine eutectic matrix Fig. 7(a) shows a TEM image of the ultrafine eutectic microstructure. The first impression is that the b and c phases exhibit a fine lamellar eutectic structure. However, when observed along another crystallographic direction, Fig. 7(b), the cross-section of the b phase is embedded in the c phase and has the shape of a particle. These observations indicate that the eutectic is actually rod-like (b phase). A statistical analysis shows that on average the dimensions of the b phase inside the eutectic matrix are about a few micrometers in length and 70–80 nm in the transverse direction. The volume fraction of the b phase in the eutectic matrix is estimated to be about 23% based on image analysis of the area fractions observed. To summarize, the
Fig. 7. Bright-field TEM images of the eutectic matrix observed along near (a) [2 0 1]b and (b) [1 1 1]b.
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eutectic matrix in the Ti60Cu14Ni12Sn4Nb10 alloy has a rod eutectic structure on an ultrafine scale, i.e., the b rods are embedded in the c matrix. The b/c interfaces have also been investigated by means of HRTEM observations. The two phases form a semicoherent interface. Fig. 8(a) and (b) shows a HRTEM image and the corresponding Fourier filtered image of the b/c interface taken along the [1 1 1]b direction. Some misfit dislocations, along with dislocations of other types, are observed at the well-bonded interface. Fig. 9(a) and (b) are the original and the Fourier filtered HRTEM images of the b phase inside the nano/ultrafine eutectic matrix along the [1 1 1]b direction. The inset is a Fourier transformed pattern of the HRTEM image, in which obviously elongated spots can be observed. In Fig. 9(b), stacks corresponding to the elongated spots in the inset are
Fig. 9. (a) Original and (b) the Fourier filtered HRTEM images of the b phase inside the nano/ultrafine eutectic matrix along the [1 1 1]b direction. The inset shows a Fourier transformed pattern of the HRTEM image.
observed; also, dislocations and lattice distortions are found in the b phase. These observations suggest that many lattice defects are stored in the as-cast eutectic matrix. The microstructural features uncovered in this section are correlated with mechanical properties in the following section.
Fig. 8. (a) HRTEM image and (b) the corresponding Fourier filtered image of the b-Ti/c-CuTi interface in the eutectic matrix.
4.2.3. Structure after deformation Fig. 10 shows the SEM image for the alloy after the compression test. During the uniaxial compression the load was applied in the horizontal direction. Some deformation bands (denoted by black arrows) are clearly seen, which are often described as shear bands [16]. The spacing of these bands is about 1–2 lm. These shear bands travel across rather long distances, propagating through dendritic arms and the eutectic colonies, roughly along the direction 45
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Fig. 10. SEM secondary electron image of the surface of the deformed alloy, showing the cross-colony shear bands (black arrows) and the intracolony slip bands (white arrows). The inset shows an SEM backscattered electron image of the region enclosed by the dashed rectangle, showing the dendrite arms and the propagation of the slip bands.
to the compression loading direction (the maximum shear stress direction). Besides these shear bands, there are shorter bands that appear to be of a high density (denoted by white arrows). Most of these bands propagate inside the
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dendrites and across the gap (matrix) in between the closely spaced dendritic arms, see the inset in Fig. 10. The spacing of these bands is less than 700 nm, and the direction of these bands has no special relationship with the compression direction, indicating that these are slip bands. By use of TEM, the major shear bands that travel long distances such as those in Fig. 10 were not observed. They are neither very large in number nor in width, making them difficult to be caught in TEM sample preparation. However, the slip bands, which we believe contribute the most to plastic deformation, were easily found. From Fig. 11(a), the spacing of the slip bands is estimated to be in the range 90–300 nm. The value is in reasonable agreement with that obtained from Fig. 10. Two slip systems were active. By a series of tilting experiments in the TEM instrument, the two systems were determined to be ð1 1 2Þ½1 1 1 and ð1 1 2Þ½1 1 1, as shown in Fig. 11(a). In a given ‘‘dendrite colony’’ (a dendrite with arms), the slip bands extend along virtually the same direction. The extensive slip deformation induced pile-ups of dislocations at the interfaces between the dendrite arm and the matrix, see Fig. 11(b) and (c). The dislocation pile-ups, as often invoked in the interpretation of the Hall–Petch relationship, lead to stress concentration and thus the transmission of the slip across the interface into the matrix, see Fig. 11(d). The steps left at the interface due to such crossing events are shown in Fig. 11(b), denoted by black arrows. The matrix between
Fig. 11. Bright-field TEM images of the deformed alloy. (a) Slip bands in each ‘‘dendrite colony’’ aligned approximately along the same direction; (b) pileup of dislocations and steps (black arrows) left by slip across the interface; (c) enlarged view of the dislocation pile-up; (d) slip band propagating into the harder matrix phase.
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the dendrite arms is mostly the c-CuTi phase, again in semicoherent relationship with the b dendrite arms. As such, the slip band runs straight across the interface, Fig. 11(d). 5. Discussion The compressive and tensile stress–strain curves of the Ti60Cu14Ni12Sn4Nb10 alloy shown in Fig. 1 demonstrate several desirable properties that can be linked to its unique microstructure. The mechanical properties of composites are sometimes interpreted using the rule of mixtures [17] P c ¼ P dV d þ P mV m
ð1Þ
where Pc is the property of the composite, and Pd and Pm are the properties of the dendrite and matrix, respectively. Vd andVm are the volume fractions of the dendrite and matrix, respectively. In our alloy, the dendrite b phase is the soft phase while the nano/ultrafine eutectic matrix is the hard phase. The volume fraction of dendrite is estimated to be approximately 50% based on image analysis. As expected from Eq. (1), the yield strength shown in Fig. 1 would be lower than that in Ref. [8], in which the volume fraction of the softer dendrite is about 20% and the yield strength of that composite is correspondingly higher (at 1.5 GPa). Obviously, the property desired can be adjusted by simply tailoring the volume fractions of the participating phases, when designing the composition of the alloy [8,18]. Even with the high volume faction of b-Ti, our alloy still exhibits a yield strength higher than commercial Ti alloys [19]. The high strength of the hard matrix comes from several microstructural features discussed in Section 4.2, including the UFG eutectic structure, the high contents of multiple solutes that lead to solution hardening, and the lattice defects in the as-cast alloy (i.e., high densities of dislocations, faults and lattice distortions revealed in TEM as described above (Fig. 9)). That the UFG eutectic is the strengthening phase can be easily appreciated by the observation that dislocations pile-up at the dendrite/matrix boundary before they enter the matrix from the dendrites. The trace amount of TiC phase, an unexpected byproduct due to contamination during sample preparation, possibly contributes as another strengthening phase. The results of Section 4.2.3 suggest that the numerous slip bands due to dislocation activities in the much larger ductile dendrites contribute to the overall ductility of this alloy. Furthermore, the accumulation of dislocations and interactions of the different slip systems contribute to the obvious strain hardening. With escalating stresses, the slip band transmits across the interfaces and propagates in the harder matrix. Without the ductile reinforcements, monolithic UFG alloys usually cannot sustain high rates of strain hardening at high stress levels, making it difficult to attain uniform tensile ductility [20]. As mentioned before, in addition to the slip bands that propagate within a dendrite colony, there are also shear bands that travel past many dendrite colonies as well as eutectic colonies. These sporadic shear bands, observed in
SEM, would also contribute to plastic strain to some extent, but not to the degree of that discussed in Section 4.2.3 for the plasticity inside the dendrite colonies. The eutectic structure, especially when decorated with a high density of defects as depicted in Fig. 9, is expected to be very strong when compared with the dendrite phase. Their contribution to the strength is expected, but their contribution to plastic strain and strain hardening is not expected to be very large, just as in the case for nano/UFG multilayers [21]. The graded length scales of the features (grain size, arm spacing, etc.) along the radial direction observed in this cast alloy can play a positive role in providing useful properties. A similar grain-size-graded structure has been achieved by means of surface mechanical attrition treatment [22], and was found to impart beneficial properties. Hanlon et al. [23] recently studied the fatigue response of nanocrystalline and UFG metals and alloys. They found that grain refinement generally leads to an increase in the resistance to failure under stress-controlled fatigue. A high-strength nanostructure is desirable at the surface where the cracks initiate. However, Hanlon et al. also concluded that nanostructures would have a generally deleterious effect on the resistance to the growth of a fatigue crack once it starts. In other words, one would prefer less refined structure in the interior in order to stop crack growth. In this context, the micro/nanostructure gradient we observed naturally offers a functionally graded material for such applications. 6. Summary A Ti60Cu14Ni12Sn4Nb10 alloy has been processed to be composed mostly of a micrometer-sized dendritic Ti-based bcc solid solution uniformly dispersed within a UFG matrix. The matrix is characterized by a rod eutectic microstructure. As the system is approximately a pseudo-binary undergoing a metastable eutectic reaction, the eutectic phases are primarily a b-Ti(Nb,Sn) solid solution alternating with a c-TiCu(Ni) intermetallic compound, with the b/c orientation relationship being [1 1 1]b//[2 2 1]c and ð110Þb ==ð110Þc . The dendrites provide the large plastic strains and strain hardening (e.g., dislocation storage) capability, whereas the strength mainly comes from the matrix where the UFG eutectic structure and the high concentrations of multiple solutes and crystallographic defects strengthen the alloy. We obtained from TEM clear and direct evidence of the microscopic dislocation mechanisms, in terms of dislocation pile-ups at the strengthening interfaces before the dislocations come from the softer phase and enter the harder phase, as well as extensive build-up of networks of slip bands in the dendrite arms. The transmission of slip across the interfaces (both dendrite/matrix and the semi-coherent boundaries between the eutectic layers) requires high stresses, hence leading to high strength. The slip propagation also produces slip steps and hence large plastic strains, and promotes work hardening due to the interactions among the accumulating dislocations and
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multiple activated slip systems. The numerous interphase boundaries in the microstructure accentuate the role of interfaces in the strengthening and ductilization of the alloy. This work demonstrates one strategy for an alloy to achieve high strength, large (tensile) ductility, high strain hardening rates and functionally graded properties, all at the same time. While it is also possible to obtain such properties in an elemental [9] or simple binary system [24], we have shown that a multi-component recipe offers more ‘‘knobs to turn’’. In this particular Ti alloy case, the addition of Nb stabilizes the b phase, allowing the ductile bcc b to be the primary phase (rather than the hexagonal close packed a). It also effectively creates a relatively deep L ! b + c metastable eutectic, promoting the formation of a UFG eutectic structure. The multiple solutes also suppress the diffusion and coarsening of the eutectic structure, and are also helpful for strengthening purposes. Our in situ composite route can also yield useful materials in bulk form via a one-step casting process.
Acknowledgments The authors acknowledge the support of the National Natural Science Foundation of China (Grant Nos. 50125103 and 50323009), the Shenyang Center for Interfacial Materials (the MANS research team), the Alexander von Humboldt Foundation, and the EU within the framework of the RTN-Network on bulk metallic glasses (HPRN-CT-2000-00033).
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References [1] Dieter GE. In: Mechanical metallurgy. Boston: McGraw-Hill; 1986. p. 212. [2] Mccandlish LE, Kear BH, Kim BK. Mater Sci Technol 1990;6:953. [3] He L, Ma E. Nanostructured Mater 1996;7:327; He L, Ma E. J Mater Res 1996;11:72. [4] Lu K, Wang JT, Wei WD. J Appl Phys 1991;69:522. [5] Jia D, Wang YM, Ramesh KT, Ma E, Zhu YT, Valiev RZ. Appl Phys Lett 2001;79:611. [6] Koch CC. Scripta Mater 2003;49:657. [7] Tellkamp VL, Melmed A, Lavernia EJ. Metall Mater Trans A 2001;32:2335. [8] Dai QL, Sun BB, Li Y, He G, Eckert J, Ma E, et al. J Mater Res 2004;19:2557. [9] Wang YM, Chen M, Zhou F, Ma E. Nature 2002;419:912. [10] He G, Eckert J, Lo¨ser W, Schultz L. Nature Mater 2003;2:33. [11] Ma E. Scripta Mater 2003;49:663. [12] Wang YM, Ma E. Acta Mater 2004;52:1699. [13] Guo FQ, Wang HJ, Poon SJ, Shiflet GJ. Appl Phys Lett 2005;86:091907. [14] Youssef KM, Scattergood RO, Murty KL, Horton JA, Koch CC. Appl Phys Lett 2005;87:091904. [15] Cheng S, Ma E, Wang YM, Kecskes LJ, Youssef KM, Koch CC, et al. Acta Mater 2005;53:1521. [16] Kim KB, Das J, Baier F, Eckert J. Appl Phys Lett 2005;86:171909. [17] Chadwick GA. Metal Sci J 1975;9:300. [18] He G, Eckert J, Lo¨ser W, Hagiwara M. Acta Mater 2004;52:3035. [19] Manero JM, Gil FJ, Planell JA. Acta Mater 2000;48:3353. [20] Ma E. Nature Mater 2003;2:7. [21] Misra A, Hirth JP, Hoagland RG, Embury JD, Kung H. Acta Mater 2004;52:2387. [22] Tao NR, Zhang HW, Lu J, Lu K. Mater Trans 2003;44:1919. [23] Hanlon T, Kwon YN, Suresh S. Scripta Mater 2003;49:675. [24] Louzguine DV, Kato H, Louzguina LV, Inoue A. J Mater Res 2004;19:3600.