Nanostructured Composites in Multicomponent Alloy Systems

Materials Transactions, Vol. 44, No. 10 (2003) pp. 1999 to 2006 Special Issue on Nano-Hetero Structures in Advanced Metallic Materials #2003 The Japan Institute of Metals

Nanostructured Composites in Multicomponent Alloy Systems Ju¨rgen Eckert1;2 , Guo He1;3 , Jayanta Das1;4 and Wolfgang Lo¨ser1 1

IFW Dresden, Institut fu¨r Metallische Werkstoffe, Postfach 270016, D-01171 Dresden, Germany Technische Universita¨t Darmstadt, FB 11 Material- und Geowissenschaften, FG Physikalische Metallkunde, Petersenstrae 23, D-64287 Darmstadt, Germany 3 Light Materials Group, National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan 4 Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur-721302, India 2

A group of novel nanostructured composites fabricated by different casting methods is presented. Nb, Ta and Mo are added into Ti-base bulk metallic glass (BMG)-forming alloys and Nb is added to the Zr-base BMG-forming alloy to induce the formation of dendrite/ nanostructured matrix composites. These composites exhibit high fracture strength of up to 2400 MPa. Both Nb-bearing Ti- and Zr-base composites exhibit over 14% plastic strain upon room temperature compression. The Ti60 Cu14 Ni12 Sn4 Nb10 composite also exhibits over 7% room temperature tensile plastic strain. The high strength of the composites is attributed to the nanostructured matrix. The large plasticity is due to the retardation of excessive localized shear banding in the matrix by the presence of the ductile dendrites. The significant work hardening before fracture is attributed to the deformation behavior of the dendritic solid solution. (Received March 24, 2003; Accepted April 22, 2003) Keywords: nanostructured materials, bulk metallic glass, composites, mechanical properties, deformation, fracture

1.

Introduction

In the past decade, bulk metallic glass (BMG) and bulk nanostructured metallic composites have been comprehensively investigated regarding their composition, fabrication methods, thermodynamic behavior as well as their mechanical properties. These two kinds of materials are very often mentioned together because of two reasons: (1) they exhibit very similar deformation and fracture behavior; and (2) bulk nanostructured materials can be fabricated from modified multicomponent BMG alloys. Due to their high yield strength, high hardness and unique physical and chemical properties, BMGs and bulk nanostructured materials are considered as very promising advanced engineering structural materials.1–3) Among these materials, particularly Zrbase BMGs are widely studied and their mechanical behavior is well understood because they exhibit very good glassforming ability and can be easily cast into large samples.4,5) Like conventional amorphous alloys, Zr-base BMGs present brittle room temperature deformation behavior under loading.4–16) This shortcoming is due to the formation of highly localized shear bands during deformation. Further deformation leads to softening in the localized shear bands. Such inhomogeneous deformation behavior yields that the final fracture occurs along the planes of those softened shear bands leading to catastrophic failure of the material with little overall plastic deformation in an apparently brittle manner. These deformation and fracture properties are very similar to that of nanostructured metallic materials.3,17,18) Generally, nanostructured materials are often extremely hard and brittle. Almost no ductility appears in tension for grain sizes smaller than about 25 nm.3) The low ductility in tension is presumably attributable to the mechanical instability due to lack of strain hardening. For BMGs, three methods to prevent inhomogeneous deformation have been introduced so far, which are particle reinforcement (adding particles into the alloy during processing19,20) and in situ formed particles21)), fiber reinforce-

ment (continuous fibers22–24) and random short fibers25)) and in situ formed ductile phase precipitates.26–28) The latter method can significantly enhance the room temperature ductility of BMGs by introducing the ductile phase, which usually has a -Ti-type cubic structure and dendritic morphology. Recently, this method was also used to enhance the ductility of bulk nanostructured materials in our group.29) The in situ formed ductile dendrites dispersed in a nanostructured matrix can act as obstacles restricting the excessive deformation of the matrix by isolating the highly localized shear bands in small and discrete inter-dendritic regions.29) The dendrites can also contribute to the plasticity by dislocation initiation and propagation. The successful fabrication of the dendrite-containing Ti-Cu-Ni-Sn-Ta(Nb) nanostructured composites is encouraging to further explore the universal aspects of such nanostructured composites in multicomponent alloy systems. In this paper, the design of new nanostructured composites, the fabrication methods and some results on microstructure and mechanical behavior will be presented. 2.

Experimental Procedure

2.1 Design of nanostructured composites We have worked on both Zr-base and Ti-base multicomponent alloy systems because Zr-base alloys exhibit a very high glass-forming ability and nanostructures are easy to achieve; Ti-base alloys are very attractive for light-weight applications. The alloys were designed by modification of well known BMG compositions. In many cases, Zr-base BMGs with addition of elements, e.g., V, Cr, Mo, Fe and Co, tend to form intermetallic compounds due to the decrease in glass-forming ability.30) Ti-base BMGs have the same problem. In fact, the potential elements for forming a dendritic bcc-phase must have a tendency to form a solid solution with Zr or Ti and must possess a high melting temperature, in order to induce the formation of primary dendritic crystals at high temperatures. Possible candidates

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J. Eckert, G. He, J. Das and W. Lo¨ser

are Nb, Ta and Hf for Zr-base BMGs and Nb, Ta and Mo for Ti-base BMGs because these elements are -phase formers and stabilizers in Zr or Ti alloys according to the binary phase diagrams.31) They can infinitely be dissolved into Zr or Ti to form a bcc--type solid solution in the range of 1500-2200 K. BMGs usually have lower melting temperatures (about 9001400 K). A proper composition combined with optimized processing can form a primary -phase upon cooling from the melt, and the remaining liquid then transforms into a nanostructure at relative low temperatures due to its highly dense random packed structure.32) As a result, the primarily solidified -phase exhibits a dendritic morphology, and is homogeneously distributed in the matrix. The proper composition range for forming a dendritic -phase embedded in a nanostructured matrix is narrow because the alloys have more than five constituents between which intermetallics can easily form, which counteracts the reinforcement by the dendritic -phase and deteriorates the ductility of the composites. Thus, suppression of intermetallic compounds is very important. 2.2 Fabrication methods The samples were prepared in several steps. First, pure Zr or Ti and the refractory metal were arc-melted to produce an intermediate supersaturated binary alloy. The master alloys were prepared by arc-melting the intermediate alloy together with the mixtures of the other metals. Then cast cylinders with 50-75 mm in length and various diameters (2-5 mm) were prepared by suction casting, centrifugal casting and melting injection casting into copper molds, respectively. Alternatively, levitation melting with a cold-crucible casting device was used to cast samples with a length of 100 mm and a diameter of 10 mm. The oxygen content of the as-cast cylinders prepared by the different methods was chemically analyzed to be less than 200 ppm. 2.3

Microstructure characterization and mechanical tests The microstructures and the phases of the as-cast samples were characterized by using a JEOL-JSM6400 scanning electron microscope (SEM) with an electron microprobe analysis unit and a JEOL-2000FX electron transmission microscope (TEM) operated at 200 kV accelerating voltage coupled with energy-dispersive X-ray analysis (EDX), as well as by using a Siemens D 5000 X-ray diffraction (XRD) facility with CuK radiation. The deformation behavior and the mechanical properties were investigated by using an Instron 8562 testing machine at a strain rate of 1 104 s1 at room temperature. The test samples were prepared according to ASTM standards. The deformation and fracture observations were carried out by SEM. 3.

Results

3.1 As-cast microstructure of different alloy systems For different alloy systems, the composite microstructures are quite different in their morphologies, sizes and homogeneity. Mo can improve the formation of the dendritic -phase and can refine the microstructure leading to very high fracture strength. Figure 1(a) shows the as-cast microstruc-

ture of a Ti47:5 Cu21:85 Ni19 Sn6:65 Mo5 alloy. Three phases can be identified, i.e., -Ti solid solution, a Ti- and Sn-enriched bright phase and Ti2 Ni black particles.33) For the Ti53 Zr6 Cu26:5 Ni4:5 Ta10 alloy, 10 at% Ta induce well developed dendrites dispersed in the matrix as shown in Fig. 1(b). EDX and XRD analysis indicate that the dendritic phase is a -Ti(Ta) solid solution. The residual Ti together with the other constituents forms the fine-grained matrix phases. Addition of 10 at% Nb can induce a perfect dendritic Ti(Nb,Sn) phase dispersed in a nanostructured matrix in the Ti60 Cu14 Ni12 Sn4 Nb10 , alloy.29) Figure 1(c) shows its equiaxed dendritic morphology. The appropriate size and volume fraction of the dendritic phase significantly improves the ductility of bulk nanostructured materials.29) For Zr66:4 Cu10:5 Ni8:7 Al8 Nb6:4 , suction casting was used for the preparation of the cast cylinders. Finer flower-like dendrites (Fig. 1(d)) dispersed a nanocrystalline matrix grain sizes in the range of 50 to 150 nm have been obtained compared to those previously found for injection cast samples.34) With increasing Zr and Nb contents, a mixture of coarse and fine dendrites dispersed in a matrix with grain sizes of about 200 to 400 nm is formed in the Zr73:5 Cu7 Ni1 Al9:5 Nb9 alloy. The volume fraction of the dendritic phase is more than 70 vol%, as shown in Fig. 1(e). Using levitation melting and coldcrucible casting, a coarser dendritic morphology is obtained in a 10 mm diameter cylinder (Fig. 1(f)). The volume fraction of the dendritic phase is more than 90 vol% evaluated from the micrograph. Compared to the suction cast 5 mm diameter cylinders, the cooling rate during solidification is much lower for the 10 mm diameter cylinders. The lower cooling rate promotes coarsening of the dendrites and leads to a larger volume fraction of the dendritic phase. This, in turn, affects the mechanical behavior of the material. Figure 2 displays bright-field TEM images of the as-cast microstructure of Ti60 Cu14 Ni12 Sn4 Nb10 . Figure 2(a) shows the dendritic phase dispersed in the matrix. The inset in Fig. 2(a) shows the selected-area diffraction pattern taken along the [1 11] zone axis of one of the dendrites, which further confirms that the dendritic phase is bcc--Ti(Nb,Sn). The more detailed microstructure of the matrix is shown in Fig. 2(b), from which a very finely layered microstructure (about 10-50 nm) can be distinguished. Some moire´ fringe structures are often displayed in this kind of nanostructured composites.29) A selected-area diffraction pattern (inset in Fig. 2(b)) taken from the matrix corroborates the nanostructure. The average composition of the dendrites and the matrix is determined by EDX analysis to be Ti61:3 Cu4:5 Ni3:1 Sn5:0 Nb26:1 and Ti55:4 Cu19:5 Ni19:5 Sn0:0 Nb5:5 , respectively. Figure 3 shows a bright-field TEM image of an as-cast Zr73:5 Cu7 Ni1 Al9:5 Nb9 10 mm diameter cylinder. The selected-area diffraction analysis confirms that the dendritic phase is bcc--Zr(Nb) (the inset in Fig. 3(a) shows a SAD pattern taken from the dendritic phase). EDX analysis reveals an average dendrite composition close to Zr80 Cu3:5 Ni0:5 Al7 Nb9 . Figure 3(b) shows a detail of the matrix and a selected-area diffraction pattern (inset in Fig. 3(b)) taken from the matrix corroborating the nanostructure. The average composition of the matrix is determined by EDX analysis to be Zr64:5 Cu20 Ni3:5 Al10 Nb2 .

Nanostructured Composites in Multicomponent Alloy Systems

Fig. 1 SEM backscattered electron images of as-cast microstructures. (a) alloy A (injection casting); (b) alloy B (injection casting); (c) alloy C (injection casting); (d) alloy D (suction casting); (e) and (f) alloy E (suction casting and cold-crucible casting, respectively).

Fig. 2 Bright-field TEM images of alloy C. (a) showing the -Ti(Nb,Sn) dendrites; the inset is the selected-area diffraction pattern taken along the [1 11] zone axis of one of the dendrites. (b) showing the nanostructured matrix; the inset displays the selected-area diffraction pattern taken from the matrix.

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Fig. 3 Bright-field TEM images of alloy E prepared by levitation melting and cold-crucible casting. (a) showing the -Zr(Nb) dendrites; the inset is the selected-area diffraction pattern taken along the [1 11] zone axis of one of the dendrites. (b) showing the nanostructured matrix; the inset displays the selected-area diffraction pattern taken from the matrix.

Table 1 Summary of the compressive test data: Young’s modulus E, yield stress y , strain at the yield point "y , ultimate compression stress max , and plastic strain "p . No

Composition/at%

Sample

E/

y / "y / max / "p /

GPa GPa % MPa %

A Ti47:5 Cu21:85 Ni19 Sn6:65 Mo5 1)

3 mm 106 2150 2.6 2246 0.5

B Ti53 Zr6 Cu26:5 Ni4:5 Ta10 1)

3 mm 130 1852 2.8 1857 0.6

C Ti60 Cu14 Ni12 Sn4 Nb10 1)

3 mm 66

1340 2.2 2401 14.6

D Zr66:4 Cu10:5 Ni8:7 Al8 Nb6:4 2)

5 mm 93

1745 2.0 1922 0.6

Zr66:4 Cu10:5 Ni8:7 Al8 Nb6:4 5)

5 mm 72

1638 2.4 1800 1.3

E1 Zr73:5 Cu7 Ni1 Al9:5 Nb9 2)

5 mm 89

1555 1.9 1740 2.1

E2 Zr73:5 Cu7 Ni1 Al9:5 Nb9 3) E3 Zr73:5 Cu7 Ni1 Al9:5 Nb9 4)

5 mm 84 10 mm 82

1305 1.8 1771 8.9 1231 1.7 1754 15.7

Fig. 4 Cast-in pores in as-cast alloy C. The arrows indicate pores formed in the inter-dendritic region.

1) Injection casting; 2) suction casting; 3) centrifugal casting; 4) levitation melting and cold-crucible casting; 5) data of Ref. 28

3.2 Casting defects For the present multicomponent alloy systems, almost all alloys exhibit very large temperature intervals between the liquidus and solidus temperatures. The primary dendrites solidify at high temperatures, e.g. at around 1500-2000 K for -Ti(M) or -Zr(M) solid solutions (M = refractory metal). The remaining liquid then solidifies at around 1000 K. The large liquid shrinkage (including thermal shrinkage and liquid/solid phase transformation shrinkage) often leads to cast-in pores dispersed in the inter-dendritic regions. Figure 4 shows cast-in pores in the inter-dendritic regions observed for as-cast Ti60 Cu14 Ni12 Sn4 Nb10 cylinders. Dendrites surrounded by the last solidified liquid emerge from the matrix in Fig. 4. The pores in the inter-dendritic regions show irregular shape and distribution, which must significantly affect the mechanical properties of the materials. It is noticed that such kind of cast-in pores is very often formed in the refractory metal-bearing nanostructured multicomponent alloys. To eliminate these cast defects, some modified casting methods, e.g., pressure casting, semisolid casting, etc., should be used. Further optimized alloy design and microstructure control are also needed.

3.3 Mechanical properties Table 1 summarizes the mechanical properties of the different alloys. The Mo-bearing alloy A exhibits a high yield strength of 2150 MPa together with a Young’s modulus of 106 GPa and over 2% elastic strain. It is a highly elastic material with an elastic energy of about 24 MJ/m3 .35) However, the material exhibits low ductility of only 0.5% plastic strain before fracture. The Ta-bearing alloy B exhibits 1852 MPa yield strength and 130 GPa Young’s modulus, but also low plastic strain. Both Ti-base alloys A and B are strong but exhibit low ductility. The Nb-bearing alloy C, however, is less strong but displays a very good ductility. It exhibits 1340 MPa yield strength and 14.6% plastic strain accompanied by significant work-hardening. The low Young’s modulus of 66 GPa indicates the high elastic capacity of the material.35) Tensile tests also show a very good ductility of this alloy, i.e., the room temperature tensile plastic strain reaches 7%.36) The suction cast Zr-base alloy D exhibits 1745 MPa yield strength, 93 GPa Young’s modulus and 0.6% plastic strain before fracture. Compared to injection cast specimens of the same alloy with dendrites dispersed in a glassy matrix


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2600 2400

C

A

2200 2000

D

Stress / MPa

1800

B

E3

E2

E1

1600 1400 1200 1000 800 600 400 200 0 0

2

4

6

8

10

12

14

16

18

Strain / %

Fig. 5 Room temperature compressive stress-strain curves of as-cast alloys A through E.

(y ¼ 1638 MPa, E ¼ 72 GPa, and "p ¼ 1:3%),28) the present sample with a nanostructured matrix exhibits higher strength and Young’s modulus but lower ductility. With high Zr- and Nb-contents, alloy E exhibits lower yield strength and low Young’s modulus but larger plastic strain. The as-cast cylinders prepared by suction casting, centrifugal casting and cold-crucible casting, respectively, exhibit different mechanical properties due to different cooling rates during solidification, which affect the size and volume fraction of the dendritic phase. Low cooling rates lead to more and coarser dendrities resulting in lower strength but higher plastic strain. The room temperature compressive stressstrain curves for all as-cast alloys A through E are shown in Fig. 5. 3.4 Deformation and fracture features Like BMGs, nanostructured metallic materials follow the shear-banding mechanism upon room temperature deformation.3,17,18) For the present nanostructured composites, shear bands can be observed on the deformed samples. Figure 6 shows a typical deformed sample surface of the Ti66:1 Cu8 Ni4:8 Sn7:2 Nb13:9 alloy on which shear bands are marked by white arrows. The rough direction of the shear bands makes an angle with the stress axis. Due to the existence of dendrites, almost all the shear bands are serpentine. This indicates that the spread of the shear bands is blocked and localized excessive shear banding can be avoided or retarded by the dendrites. The crack crossing (not along) the shear bands (indicated by black arrows in Fig. 6(b)) again reveals that excessive shear banding (which can induce cracks along the shear bands) is hindered by the dendrites. Although large plastic strain occurs before fracture, the samples of the present nanostructured composites exhibit a shear fracture16,37) and produce an almost smooth fracture surface. Observation of the details of the fracture surface shows some different features for these composite materials. For the strong Ti-base alloy A, both network-like and cleavage-like fracture surfaces can be observed (Fig. 7(a)). The network-like surface is related to the dendrite-induced ductile fracture. The cleavage-like surface is related to brittle fracture, which is induced by the nanocrystalline matrix. For the Ta-bearing alloy B, the elongated dendrites induce a

Fig. 6 Image of the deformed sample surface of Ti66:1 Cu8 Ni4:8 Sn7:2 Nb13:9 alloy. (a) is a overview of the fractured sample surface showing the relationship of stress axis, shear bands and fracture surface. (b) reveals the detail of the deformed sample surface showing shear bands. Small white arrows indicate the shear bands; black arrows indicate a crack cross to the shear bands.

stick-like fracture surface (Fig. 7(b)). The dendrite-related ‘sticks’ are expected to contribute to the ductility of the material. On the other hand, a lot of cast-in pores can be found in the inter-dendritic regions, which may be responsible for the observed low ductility of this alloy. For alloy C, a homogeneous network-like fracture surface can be observed, as shown in Fig. 7(c). The enlarged detail of the fracture surface (inset in Fig. 7(c)) shows obvious viscous flow, which is very similar to what is known for BMGs. The difference is that the viscous flow is restricted to the interdendritic region for the present nanostructured composite. This can avoid excessive localized deformation and retard the failure of the material. The dendrites not only restrict the excessive localized flow, but also contribute to the deformation by dislocation slip. The significant work-hardening found for this alloy is expected to stem from the dendritic Ti(Nb,Sn) solid solution. For the cold-crucible cast 10 mm diameter samples, a very similar deformation mechanism is observed. The fracture surface shows a rough but viscous feature (Fig. 7(d)). The large plasticity arises from the large volume fraction (over 90 vol%) of ductile dendrites. 4.

Discussion

4.1 Formation of nanostructured composites The dendrite/nanostructured matrix composites can only form under the conditions of appropriate constituent configuration of the alloy and well-controlled solidification. The

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Fig. 7 Fractography of as-cast samples. (a) alloy A; (b) alloy B; (c) alloy C, the inset shows the detail of the fracture surface; and (d) alloy E prepared by levitation melting and cold-crucible casting.

appropriate constituents are arranged to achieve the formation of the dendritic bcc--phase and, on the other hand, to maintain the composition of the residual liquid close to eutectic and having a highly dense random packed liquid structure32) so as to achieve a nanocrystalline transformation upon continuous cooling. Besides, the composition arrangement is also required to avoid or restrict the formation of intermetallic compounds to a very small volume fraction. Otherwise, the intermetallic phases induce an intrinsic brittleness of the composites that can completely counteract the ductile dendritic phase reinforcement. Beyond the composition adjustment, increasing cooling rate can also avoid the formation of intermetallic compounds. In our study, it has been found that Nb and Ta are very efficient to form nanostructured composites. They can form an appropriate size, morphology and homogeneous distribution of ductile dendrites in a nanostructured matrix for both Ti- and Zr-base multicomponent alloys. However, Ta was found to be very difficult to mix with the other constituents and very often induces cast-in pores in the inter-dendritic regions due to its high melting temperature.37) Mo can easily induce the precipitation of intermetallic compounds and the formation of cast-in pores in Ti-base composites.33) Nb can also induce the precipitation of intermetallic phases in Zr-base composites, which can be avoided by increasing the cooling rate. This may lead to the formation of a dendrite/glassy matrix composite.28) The cast-in pores in the inter-dendritic region are very often found in the present nanostructured composites because of their very large solidification temperature intervals. This

drawback can be overcome by improving the casting procedure or by using advanced casting methods such as pressure casting, semisolid casting, directional solidification, etc. Since the tensile strength of the materials is very sensitive to the presence of cast-in pores, avoiding such casting defects is expected to improve the mechanical properties of these composites. Our recent experiments reveal that pore-free Ti60 Cu14 Ni12 Sn4 Nb10 composite specimens exhibit over 7% plastic strain during tensile testing.36) 4.2 Dendrite reinforcement mechanism The nanostructured composites are usually composed of a ductile dendritic phase, a nanostructured matrix, and a few other crystalline phases, such as intermetallics. Any large size (e.g. 0.1-10 mm) brittle phase in the microstructure is harmful for the ductility of the material. The larger the volume fraction of such brittle phases, the more brittle the material becomes. Thus, for the nanostructured composites, brittle phases should be controlled to a small volume fraction. It is well known that nanostructured metallic materials are very hard but brittle.3) Therefore, the nanostructured matrix in the composites contributes to the high strength of the material. The bcc solid solution phase dendrites are ‘soft’ because there are several slip systems that can be easily actuated under stress. When the stress reaches the yield strength of the nanostructured matrix, deformation firstly takes place in the matrix by shear banding. The shear bands spreading in the matrix must interact with the micrometersize dendrites. Thus, the propagation of shear bands must follow one of the three possible mechanisms, i.e., (1) shear


bands are stopped, (2) shear bands move around the dendrites; and (3) shear bands go through the dendrites and transfer the shear deformation into the dendrites. Any of these three process will retard the further propagation of the shear bands, which contributes to reducing the localized excessive shear banding and is beneficial for achieving homogeneous deformation. Both the ‘move around’ and ‘go through’ mechanisms are expected at the beginning of yielding. The ‘go through’ mechanism is dominant at a late stage of yielding since significant work hardening occurs. The deformation behavior of the dendrites affects and contributes to the deformation behavior of the composites at the yielding stage. Based on the observations of the fracture surfaces, it is evidently shown that the matrix exhibits viscous flow before fracture. The viscous flow is related to the localized shear banding. However, this viscous flow behavior is confined to small regions (e.g. interdendritic regions). Thus, it can not induce catastrophic failure of the material. If there are large size brittle crystals or cast-in pores in the microstructure, they will wreck the connection between dendrites and nanostructured matrix leading to brittleness. It is expected that if we can the desired completely dual-phase microstructure (dendrite and nanostructured matrix) without any other crystalline phase and without casting defects, high strength and toughness can be for achieved. 5.

Conclusions

The combination of refractory metals and BMG-forming alloys can yield micrometer-sized ductile dendrite/nanostructured matrix composites by using appropriate composition control and preparation procedures. By adding Nb, Ta and Mo into Ti-base BMGs and Nb and Ta into Zr-base BMGs, dendrite/nanostructured composites have been successfully fabricated. Among them, Nb is most promising for forming defect-free composite microstructures without harmful brittle phases in order to achieve very strong and very ductile nanostructured composites. The dendrite/nanostructure composites form upon solidification during which the dendritic phase primarily solidifies and the residual liquid transforms into a nanostructured matrix. The composition of the remaining liquid and the cooling rate influence the solidification process of the matrix. An appropriate composition and a high cooling rate can avoid the precipitation of intermetallic compounds. Due to the large solidification temperature intervals, cast-in pores are often formed in the inter-dendritic region, which can degrade the strength and the ductility of composites. These defects can be avoided by using advanced casting methods and wellcontrolled casting procedures. The deformation of the dendrite/nanostructure composites follows a shear banding mechanism. However, the shear bands are restricted to small regions. This can prevent excessive localized shear banding and avoids the catastrophic failure of the composites. The dendrites retard the shear banding by ‘moving around’ and ‘go through’ mechanisms, and contribute to the plastic deformation and the significant work hardening of the materials. The nanostructured matrix contributes to the high strength.

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Acknowledgements The authors thank M. Frey, H. Grahl, M. Gru¨ndlich, A. Gu¨th, H.-J. Klau, U. Ku¨hn, S. Mu¨ller-Litvanyi, and L. Schultz for technical assistance and stimulating discussions. Funding by the EU within the framework of the RTNnetwork on bulk metallic glasses (HPRN-CT-2000-00033) is gratefully acknowledged. G. He and J. Das are very grateful for the financial support of the Alexander-von-Humboldt Foundation and the German Academic Exchange service (DAAD), respectively, for providing fellowships for their stay at the IFW Dresden. This workshop was sponsored by Nanotechnology Research Network Center of Japan and Tohoku University Materials Research Center. REFERENCES 1) A. Inoue: Acta Mater. 48 (2000) 279-306. 2) W. L. Johnson: MRS Bull. 24 (1999) 42-56. 3) C. C. Koch, D. G. Morris, K. Lu and A. Inoue: MRS Bull. 24 (1999) 5458. 4) A. Inoue, T. Zhang and T. Masumoto: Mater. Trans., JIM 30 (1989) 965-972. 5) A. Peker and W. L. Johnson: Appl. Phys. Lett. 63 (1993) 2342-2344. 6) P. E. Donovan: Acta. Metall. 37 (1989) 445-453. 7) C. T. Liu, L. Heatherly, D. S. Easton, C. A. Carmichael, J. H. Schneibel, C. H. Chen, J. L. Wright, M. H. Yoo, J. A. Horton and A. Inoue: Metall. Mater. Trans. A29 (1998) 1811-1820. 8) G. He, Z. Bian and G. L. Chen: Mater. Sci. Eng. A270 (1999) 291-298. 9) P. Lowhaphandu, S. L. Montgomery and J. J. Lewandowski: Scr. Mater. 41 (1999) 19-24. 10) A. Leonhard, M. Heilmaier and J. Eckert: Scr. Mater. 43 (2000) 459464. 11) W. J. Wright, R. Saha and W. D. Nix: Mater. Trans. 42 (2001) 642-649. 12) G. He, J. Lu, Z. Bian, D. J. Chen, G. L. Chen, G. C. Tu and G. J. Chen: Mater. Trans. 42 (2001) 356-364. 13) G. He, Z. Bian, G. L. Chen, J. Lu, D. J. Chen, G. C. Tu, G. J. Chen and X. J. Hu: J. Mater. Sci. Tech. 17 (2001) 389-398. 14) Z. Bian, G. L. Chen, G. He and X. D. Hui: Mater. Sci. Eng. A316 (2001) 135-144. 15) Z. Bian, G. He and G. L. Chen: Scr. Mater. 46 (2002) 407-412. 16) Z. F. Zhang, J. Eckert and L. Schultz: Acta Mater. 51 (2003) 11671179. 17) V. Y. Gertsman, R. Z. Valiev, N. A. Akhmadeev and O. V. Mishin: Mater. Sci. Forum 225-227 (1996) 739-744. 18) P. G. Sanders, C. J. Youngdahl and J. R. Weertman: Mater. Sci. Eng. A234-236 (1997) 77-82. 19) H. Choi-Yim, R. Busch, U. Ko¨ster and W. L. Johnson: Acta Mater. 47 (1999) 2455-2462. 20) J. Eckert, A. Kuebler and L. Schultz: J. Appl. Phys. 85 (1999) 71127119. 21) C. Fan, A. Takeuchi and A. Inoue: Mater. Trans., JIM 40 (1999) 42-51. 22) R. B. Dandliker, R. D. Conner and W. L. Johnson: J. Mater. Res. 13 (1998) 2896-2901. 23) C. P. Kim, R. Busch, A. Masuhr, H. Choi-Yim and W. L. Johnson: Appl. Phys. Lett. 79 (2001) 1456-1458. 24) R. D. Conner, R. B. Dandliker and W. L. Johnson: Acta Mater. 46 (1998) 6089-6102. 25) H. Choi-Yim and W. L. Johnson: Appl. Phys. Lett. 71 (1997) 38083810. 26) C. C. Hays, C. P. Kim and W. L. Johnson: Phys. Rev. Lett. 84 (2000) 2901-2904. 27) F. Szuecs, C. P. Kim and W. L. Johnson: Acta Mater. 49 (2001) 15071513. 28) U. Ku¨hn, J. Eckert, N. Mattern and L. Schultz: Appl. Phys. Lett. 80 (2002) 2478-2480. 29) G. He, J. Eckert, W. Lo¨ser and L. Schultz: Nature Mater. 2 (2003) 3337.

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