In situ tensile testing in SEM of Al-[Al.sub.4][C.sub.3] nanomaterials/Al-[Al.sub.4][C.sub.3] nanomaterjalide in situ SEM-tombeteim.

Besterci, Michal ; Velgosova, Oksana ; Ivan, Jozef 等

1. INTRODUCTION

The method of in situ tensile testing in SEM is suitable for investigations of fracture mechanisms because it enables one to observe and document deformation processes directly, thank to which the initiation and development of plastic deformations and fracture can be reliably described.

In our previous papers [1-8], based on papers [9-11], we used in situ tensile testing in SEM to analyse deformation processes in various types of Cu and Al based composites. In [9-11] Al-Si-Fe and Al-Si systems were studied by in situ tensile testing in SEM. The result was a design of several models of damage, which considered physical and geometrical parameters of the matrix and particles.

The dispersion-strengthened alloys Al-[Al.sub.4][C.sub.3], manufactured by mechanical alloying using powder metallurgy technology, are promising structural materials. One microstructured material of such type with 4 vol% [Al.sub.4][C.sub.3] was transformed by the equal channel angular pressing (ECAP) method in two passes into a nanocomposite material. The experimental material was pressed through two perpendicular to one another channels of a special die by route "C". The ECAP technology allows obtaining very fine-grained microstructure--nanostructure--by multiple pressings through the die. The mechanical tests of this nanocomposite material in comparison with the not extruded materials show a tensile strength increase of 40% and yield strength increase of 30%. The elongation of extruded nanomaterials decreases in average by 45%.

The aim of this paper is to analyse the fracture mechanism of the Al-[Al.sub.4][C.sub.3] nanocomposite system and to propose a damage model.

2. EXPERIMENTAL MATERIALS AND METHODS

The starting experimental materials were prepared by mechanical alloying. The Al powder of powder particle size <50 [micro]m was dry milled in an attritor for 120 min with the addition of graphite KS 2.5 thus creating 4 vol% of [Al.sub.4][C.sub.3]. The specimens were then cold pressed using a load of 600 MPa. The specimens had cylindrical shape. Subsequent heat treatment at 550[degrees]C for 10 h induced chemical reaction 4A1 + 3C [right arrow] [Al.sub.4][C.sub.3]. The cylinders were then hot extruded at 600[degrees]C with 94% reduction of the cross section. Using such a treatment temperature and time of treatment, the complete transformation of graphite takes place [11]. The extent to which C was converted to [Al.sub.4][C.sub.3] dispersoids was analysed by gas chromatographic measurements in [12]. Due to a high affinity of Al to [O.sub.2], the system also contains a small amount of [Al.sub.2][O.sub.3] particles. The volume fraction of starting [Al.sub.2][O.sub.3] was low, 1 vol%. Detailed technology description is given in [13-16].

This material with dimensions of [empty set] 10 x 70 mm was deformed by the ECAP technique in two passes at room temperature in a hydraulic press in the equipment described in [17]. In [18], a dislocation model of microcrystalline system with hard nanoparticles has been suggested. Analysis of the evolution of the microstructure and mechanical properties in ECAP is given in [19-23].

For the purposes of investigation, very small flat tensile test pieces (7 x 3 mm, gauge length is 7 mm) with 0.15 mm thickness were prepared by electroerosive machining, keeping the loading direction identical to the direction of extrusion. The specimens were ground and polished down to a thickness of approximately 0.1 mm. Finally, the specimens were finely polished on both sides by ion gunning. The test pieces were fitted into special deformation grips inside the scanning electron microscope JEM 100 C, which enables direct observation and measurement of the deformation by ASID-4D equipment. From every system five samples were prepared.

3. RESULTS AND DISCUSSION

The microstructure of the starting material with 4 vol% [Al.sub.4][C.sub.3] was fine-grained (the mean matrix grain size was 0.35 [micro]m), heterogeneous, with [Al.sub.4][C.sub.3] particles distributed in parallel rows as a consequence of extrusion. Beside the phase [Al.sub.4][C.sub.3], the systems contained also [Al.sub.2][O.sub.3] phase (about 1 vol%). Essentially, it was the remnant of oxide shells of the original matrix powder and/or shells, formed during the reaction milling in attritors.

When describing microstructures, one has to consider geometrical and morphological factors. According to the microstructure observations, the particles in our materials can be divided into three distinctive groups: A--small [Al.sub.4][C.sub.3] particles, identified by TEM, with mean size approximately 30 nm, which made up to 70% of the dispersoid volume fraction; B--large [Al.sub.4][C.sub.3] particles with mean size between 1 and 2 gm, identified by scanning electron microscopy and on metallographic micrographs; and C--large [Al.sub.2][O.sub.3] particles with mean size of 1 [micro]m, found on metallographic micrographs and identified by scanning electron microscopy. Morphologically, [Al.sub.4][C.sub.3] particles are elongated and [Al.sub.2][O.sub.3] particles are spherical. Let us assume that particles of all categories during the high plastic deformation are distributed in rows. Mean distance between the rows is l and between the particles h. The particles are spherical or have only a low aspect ratio, so that they can be approximated as spherical. The experimental materials were deformed at 20[degrees]C at a strain rate of 6.6 x [10.sup.-4] [s.sup.-1] in the elastic region.

The material after ECAP is on the border of nanostructured materials. The TEM micrographs (Figs. 1 and 2) showed that the mean grain size was 100-200 nm; dislocations are present in nanograins, but mostly on the boundaries. In Fig. 2 these dislocations are weakly visible due to the tilting of the specimen. The nanostructure formation takes place most probably by a three-stage mechanism, described in [22]. This model has been experimentally verified only for several specific materials, but in our case it seems to be usable. The model includes creation of the cell structure, then formation of the transitory cell nanostructure with large angle disorientation, and finally formation of nanostructured grains with size of about 100 nm. However, here one has to consider the retarding effect due to the presence of dispersoid particles.

Deformation process of the loaded layer causes fracture of large, B-type, particles in the middle of the specimen (Figs. 3 and 4), which initializes fracture path roughly perpendicular to the loading direction (Fig. 5). The fracture path is determined also by decohesion of smaller particles (type A and/or C). Since the volume fractions of [Al.sub.4][C.sub.3] and [Al.sub.2][O.sub.3] particles are small, their distribution in lines does not influence the trajectory of fracture, which has low relative deformation [epsilon] = 0.135. Unlike the microstructured Al-based composites, in this case it has been shown that the nanograin boundaries play an important role. In the final phase (Fig. 6) a crack propagates along the nanograin boundaries, which has been observed experimentally on the crack line (profile).

[FIGURE 1 OMITTED]

[FIGURE 2 OMITTED]

[FIGURE 3 OMITTED]

[FIGURE 4 OMITTED]

A detailed study of the deformation changes showed that the crack initiation was caused by decohesion and occasionally also by rupture of large particles. Decohesion is a result of different physical properties of different phases of the system. The Al matrix has significantly higher thermal expansion coefficient and lower elastic modulus (from 23.5 to 26.5 x [10.sup.-6] [K.sup.-1] and 70 GPa) than both [Al.sub.4][C.sub.3] (5 x [10.sup.-6] [K.sup.-1], and 445 GPa) and [Al.sup.2][O.sup.3] (8.3 x [10.sup.-6] [K.sup.-1], and 393 GPa), respectively.

Large differences in the thermal expansion coefficients result in high stress gradients, which arise on the interphase boundaries during the hot extrusion. Since [[alpha].sub.matrix] > [[alpha].sub.particle], high compressive stresses can be expected. However, because the stress gradients arise due to the temperature changes, during cooling (which results in the increase of the stress peaks) their partial relaxation can occur. Superposition of the external load and the internal stresses can initiate cracking at interphase boundaries. This is in accordance also with the dislocation theories, which argue that the particles in a composite may cause an increase in the dislocation density as a result of thermal strain mismatch between the ceramic particles and the matrix during preparation and/or thermal treatment [18]. In our case, the coefficient of thermal expansion of the matrix is much higher than that of the secondary particles and the resulting thermal tension may relax around the matrix-particle interfaces by emitting dislocations, whose density can be calculated according to a procedure, described in [18].

[FIGURE 5 OMITTED]

[FIGURE 6 OMITTED]

Based on the microstructure changes, observed in the process of deformation, the following model (it is not a general model but one that resulted from our experiments) of fracture mechanism is proposed (Fig. 7).

A. The microstructure in the initial state is characterized by [Al.sub.4][C.sub.3] and [Al.sub.2][O.sub.3] particles, categorized as A, B and C, whose geometric parameters (l, h and d) depend on their volume fraction (Fig. 7a).

B. With increasing tensile load, the local cracks, predominantly on specimen side surfaces, are formed by rupture of large (B) and decohesion of smaller (C and/or A) particles (Fig. 7b,c).

C. With further increasing deformation of nanocomposite materials the nanograin boundaries start to play an important role. Since the volume fraction of these boundaries is high and the size of the B and C particles is equal to the matrix grain size, crack propagates preferentially along the nanograin boundaries (Fig. 7d).

[FIGURE 7 OMITTED]

4. CONCLUSIONS

The aim of the study was to evaluate the influence of selected volume fraction of [Al.sub.4][C.sub.3] (4 vol%) and [Al.sub.2][O.sub.3] (1 vol%) particles on the fracture mechanism by means of the method in situ tensile testing in SEM.

Based on the microstructure changes, obtained in the process of deformation of the dispersion-strengthened Al-[Al.sub.4][C.sub.3] alloys, a model of fracture mechanism was proposed. With increasing tensile load the local cracks, predominantly on specimen's side surfaces, are formed by rupture of large and decohesion of smaller particles. The orientation of cracks tends to be perpendicular to the loading direction. The final rupture, i.e. interconnection of the side cracks along the loading direction, takes place at nanograin boundaries.

doi: 10.3176/eng.2009.4.01

ACKNOWLEDGEMENT

This work was supported by the Slovak National Grant Agency under the Project VEGA 2/0105/08.

Received 29 May 2009, in revised form 14 August 2009

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Michal Besterci (a), Oksana Velgosova (b), Jozef Ivan (c), Pavol Hvizdos (a), Tibor Kvackaj (d) and Priit Kulu (e)

(a) Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, 043 53 Kosice, Slovak Republic; {mbesterci, phvizdos}@imr.saske.sk

(b) Department of Non-ferrous Metals and Waste Treatment, Faculty of Metallurgy, Technical University, Letna 9/A, 04200 Kosice, Slovak Republic; oksana.velgosova@tuke.sk

(c) Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, Racianska 75, 838 12 Bratislava, Slovak Republic; ummsivan@savba.sk

(d) Department of Metals Forming, Faculty of Metallurgy, Technical University, Vysokoskolska 4, 04200 Kosice, Slovak Republic; tibor.kvackaj@tuke.sk

(e) Tallinn University of Technology, Department of Materials Technology, Ehitajate tee 5, 19086 Tallinn, Estonia; pkulu@edu.ttu.ee