'University of Idaho, Moscow, Idaho 83844-3026. 2Ufa State Aviation Technical University, Ufa, Russia 450000. 3ALCOA Technical Center, Alcoa Center, PA ...
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10. No.5. pp. 691498.1998 Elscvia soiclua? Ldd 0 1998 Acta Mdallwgica Inc.
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MICROSTRUCTURE AND MICROHARDNESS OF AN AI-Fe ALLOY SUBJECTED TO SEVERE PLASTIC DEFORMATION AND AGING ON. Senkov’, EH. Frees*, V.V. Stolyarov2, R.Z. Va1iev2, and J. Liu3
‘University of Idaho, Moscow, Idaho 83844-3026 2Ufa State Aviation Technical University, Ufa, Russia 450000 3ALCOA Technical Center, Alcoa Center, PA 15069 (Accepted February 12,1998)
Abstra(ct-A nanocrystalline structure wasproduced in an Al-11 wt.% Fe alloy with the use of the novel technique of severe plastic deformahon of ingots by torsion under high imposed pressure. This technique allows a large departure of materialsfrom equilibrium. The microstructure of the alloys was studied with the use of TEM and EDS. The severe plastic deformation led to solid solubility extension of iron in the aluminum matrix, dispersion and dissolution of second phase particles, grain size reduction into the nanometer range, and partial amorphization. Microhardness of the alloy increased substantially after the deformation due to the grain refinement and the solid solubility extension. Aging of the severe plasticallydeformed samples at 100°C led tojurther increase of the microhardness due to the decomposition of the suRersaturated solid solution and precipitation-induced hardening. 01998 Acta Metallurgica Inc. INTRODUCTION
It has been recognized that severe plastic deforanation performed at low temperatures (usually less than 0.4 T,,,) can refine the microstructure of metals and alloys to the nanometer-sized range (l-3). Recent work has also demonstrated possible changes in phase compositions and formation of me&stable phases during processing by severe plastic deformation, such as a formation of a solid solution in the immisible alloy Cu-5OIAg (4), complete dissolution of the cementite in a high carbon steel (5), and an increase in the solid solubility of iron in aluminum (6). These novel constitutional and microstructural effects can lead to enhanced physical and mechanical properties. For example, an enhancement of aging processes in a hardened commercial Al-Cu-Zn-Mg-Zr alloy due to severe torsional straining has been demonstrated (7). Aluminum-iron alloys are attractive for engine applications. Alloying of aluminum with iron can increasethe high temperature strength due to a dispersion of second-phase particles (8). Unfortunately, the equilibrium solubility of iron in the aluminum lattice is very low and even at high tempemtures it does not exceed 0.03 at.% (9), and these alloys cannot be dispersionstrengthened with the use of conventional thermal treatments. The strengthened effect can be 691
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enhanced by increasing the solid solubility extension of iron in the aluminum matrix by “far-fromequilibrium” techniques such as rapid solidification (10,l l), mechanical alloying (12- 16) or even a laser treatment (17). In the present work, the severe plastic deformation approach has been used to extend the iron solubility in aluminum and to produce after aging a nanograined dispersionstrengthened structure in an Al-Fe alloy. EXPERIMENTAL
PROCEDURES
AND MATERIALS
Acast rod of the Al- 11 wt.% Fe alloy produced by ALCOA was used as a baseline material. Disk-shape samples were cut from this rod, with a diameter of 12 mm and a thickness of 0.3 mm. The samples were strained in torsion on Bridgman anvils atroom temperature to a true logarithmic strain E = 7 under a quasi-hydrostatic pressure of 6 GPa. The deformed samples were aged at the temperature of 1OOOCin boiled water for the period of time of 15 minutes to 12 hours. Transmission (JBM 2010) and scanning (Hitachi 2000) electron microscopes and a X-ray diffractometer (Siemens 5000, Cu K&radiation, h= 1S4056E) were used for microstructure and phase composition studies. The transmission electron microscope (TBM) was equipped with an energy dispersive X-ray spectrometer (EDS) Link ISIS, which allowed chemical microanalysis of the phases. Microhardness measurements were carried out on the samples with flat polished surfaces, with the use of a Vickers device under an applied load of 0.98 N for lo-15 seconds. Ten measurements at different points of the sample were used to calculate an average value of the microhardness.
Figure 1. Microstructure of the as-cast alloy.
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RESULTS AND DISCUSSION Microstructure Evolution
The initial cast alloy had a typical dendritic type microstructure and contained a fee Al matrix phase and monoclinic Alt3Fes phase, the latter in the form of dendrites, Figure 1. The mean grain size in the aluminum matrix was 15 pm. An analysis of the XRD patterns showed that virtually no iron was in solution in the aluminum matrix, in agreement with the equilibrium Al-Fe phase diagram (9). A very fine microstructure was produced in the alloy samples by the severe plastic deformation, which could be detectable only with the use of TEM. Two phases were present in the deformed alloy: fee Al-rich and monoclinic AltsFe4. Figure 2 shows typical microstructure of the Al-rich phase in bright and dark field images, together with a selected area diffraction pattern
Figure 2. Microstructure of the as-deformed alloy: aluminum-based matrix phase: (a) bright field image, (b) dark field image, and (c) microdiffraction.
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Energy (kev) Figure 3. (a) A selected grain of the aluminum-based matrix in the alloy after severe torsional straining and (b) EDS spectrum from the grain interior. (SADP) obtained from a region 0.5 pm in diameter. The and adjoining reflections were used to obtain the dark field image. The aluminum-rich phase had a homogeneous grain structure with a mean grain size of about 100 nm. According to the selected electron diffraction pattern, Figure 2c, the grains have high angle grain boundaries and random crystallographic orientations. Figure 2a shows that some grain boundariesarepoorly defmed and the contrast within the grams is not uniform, but often changes in a complex fashion that indicates a high level of internal stresses and elastic distortions in the crystal lattice. Because this contrast is present both in grains containing dislocations and those where no dislocations are present, one can suggest that the source of these internal stresses are defects within the grain boundaries. The EDS analysis of the aluminum-rich matrix showed that it contained from 1.34 to 2.24 wt.% Fe in solid solution. Figure 3 shows a grain and the EDS spectrum taken from the region in this gram, marked with an arrow. No second phase particles were observed in the grain, although l.% wt.% Fe was detected. The second phase was homogeneously distributed in the deformed alloy as particles of less than 1 pm in size, Figure 4. The EDS analysis of the compound particles showed that they contain of 35.0 to 38.0 wt.% Fe, which is between the AlsFe and AWe compositions. An analysis of the electron diffraction patterns from these particles, Figure 4b, showed that they consisted of the monoclinic AltsFe4 phase. An increase of grain sizes in the aluminum-based matrix and the simultaneous appearance of common banded contrasts at the gram boundaries were observed during aging at KWC, Figure 5. A well developed dislocation sub-grain structure and extinction contours resulted from elastic stresses inside the grains areother features of the aged microstructureThe dark field image analysis showed a specific dotted contrast on the dislocations and grain boundaries, Figure 5b, which was probably due to very fine precipitation of the second phase. The use of different diameter electron beam at the EDS analysis showed that the concentration of iron in the solid solution gradually decreased, although the average iron concentration in the ahrminum-based grains remained almost unchanged, Figure 6, indicating the iron precipitation from the solid solution during the aging.
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Energy (keV) Figure 4. (a) Microstructure and (b) corresponding SADP of the as-deformed alloy: an AlrsFe4 particle in the aluminum-based matrix, (c) EDS spectrum from the compound particle.
Microhardness The as-cast alloy had the Vicker’s microhardness of 750 MPa. After the severe torsional straining the microhardness increased to 1750 MPa, being higher by a factor of 2.3 as compared to the cast state. Such a considerable increase in the microhardness may be caused by several reasons, including grain refinement, increasing the defect density and internal stresses, fragmentation of the second-phase particles, and formation of the supersaturated solid solution of iron in the aluminum matrix. Artificial aging of the as-deformed alloy at 100°C led to a further increase of the microhardness, and the maximum value of 3020 MPa was reached after 5 hours of aging, Figure 7. This large increase in the microhardness during aging was apparently caused by a precipitation
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Figure 5. Microstructure of the severe plastically deformed alloy after aging at 100°C for 4 hours: (a) bright field and (b) dark field images.
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Figure 6. Concentration of iron EDS measured at the eleclron-beam spots of (0) 20 nm and (0) 0.5 pm vs. aging time.
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Time (h) Figure 7. Dependence of the microhardness on aging time at 100°C.
of the second phase particles on dislocations and grain boundaries, although some relaxation processes also occurred. Further increase of the aging time, up to 12 hours, led to a continuous decrease of the microhardness, however the values were still higher than that in the as-deformed State.
The cha.nges of the microhardness during aging of the as-deformed alloyrepresent the typical behavior of hardened aluminum alloys on aging. Obviously, the first stage of aging, when the microhardness increases, is due to a redistribution of the iron in the supersaturated solid solution and precipitation of the iron on dislocations and grain boundaries. The second stage of the aging, when thiemicrohardness decreases, is obviously caused by a coarsening of the second-phase particles, grain coarsening in the matrix phase and relaxation of the internal stresses. A very important result of the present work is the possibility of aging of the conventionally nonhardenable alloy by employing severe plastic deformation to produce a supersaturated solid solution, with a considerable strengthening effect after aging. CONCLUSIONS 1. Severe plastic torsional straining of cast Al-l lwt.% Fe alloy led to microstructural refinement and formation of a nanocrystalline structure in an aluminum-based matrix with a homogeneous, distribution of small (< 1 pm) second phase particles. 2. Asupersaturated solid solution of iron in the aluminum matrix with a maximum solubility of 2.2 wt.% was obtained during the severe plastic deformation, which allowed aging (hardening) of the conventionally non-hardenable alloy to occur. 3. The ,microhardness of the as-cast alloy of 750 MPa was increased up to 1750 MPa after the severe plastic deformation. The subsequent artificial aging at 100°C led to a further increase in the microh,ardness up to 3020 MPa due to decomposition of the supersaturated solid solution and precipitation strengthening.
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REFERENCES 1. 2. 3. 4. 5. 6.
7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17.
Valiev, R.Z., Korznikov, A.V. and Mulykov, R.R., Materials Science and Engineering A, 1993, Al68 (2), 141. Valiev, R.Z., ed., Ultrafine-Grained Materials Prepared by Severe Plastic Deformation. Special Issue, Annales de Chimie-Mater. Science, Fr. 5-6, 1996. Pavlov, V.A., Physics of Metals and Metallography, 1989,67,924. Shen, H.. Li, Z., Gunther, B., Korznikov, A.V. and Valiev, R.Z., Nanastructructured Materials, 1995,6 (l-4), 385. Ivanisenko, Yu.V., Korznikov, A.V., Safarov, I.M., Valiev, R.Z., Nanostructured Materials, 1995, 6 (l-4), 433. Senkov, O.N., Froes, F.H., Stolyarov, V.V., Valiev, R.Z. and Liu, J., in Advanced Particulate Materials and Processes 1997. eds. EH. Fmes and J.C. Hebeisen, Metal Powder Industries Federation, Princeton, NJ, 1997, p. 95. Stolyarov, V.V., Latysh, V.V., Salimonenko, VA., Islamgaliev, R.K., Valiev, R.Z., Materials Science and Engineering A, 1997, A234-236,339. Kim, Y.W. and Griffith, W.M., eds., Dispersion StrengthenedAluminum Alloys, TMS, Warrendale, PA, 1988. Kattner, U.R., in Binary Alloy Phase Diagrams, ed. T.B. Massalski, ASM, Metals Park, OH, Vol. 1, 1986, p. 147. Tonejc, A. and Bonefacic, A., Journal ofApplied Physics, 1968,40 (l), 419. Young, R.M.K. and Clyne, T.W., Scripta Metallurgica et Materialia,1981, 15(11), 1211. Mukhopadhyay, D.K., Suryanarayana, C. and Froes, F.H., Metallurgical Materials Transactions A, 1995,26A (8), 1939.. Fadeeva, V.I. and Leonov, A.V., Materials Science Forum, 1992,88-90,481. Huang, B., Tokizane, N., Ishihara, K.N., Shingu, PH. and Nasu, S., Journal of Non-Crystalline Solids, 1990, 117/118,688. Mukhopadhyay, D.K., Suryanarayana, C. and Froes, F.H., Scripta Metallurgica et Materialia, 1994,31 (3), 333. &ecoglu,M.L., Suryanarayana, C. andNix, W.D.,MetallurgicalMaterials TransactionsA, 19%, 27A (4), 1033. Uglov, A.A., Ignatev, M.B., Titov, VI., Soviet Physics Doclady, 1990.35 (ll), 977.
