Deformation behaviour and microstructure-mechanical-property correlations for RR58 aluminium alloy

Journal of Materials Processing Technology, 29 (1992) 91-101

91

Elsevier

Deformation behaviour and microstructuremechanical-property correlations for RR58 aluminium alloy K.P. Rao Department of Manufacturing Engineering, City Polytechnic of Hong Kong, Kowloon, Hong Kong K. Sivaram Central Metal Forming Institute, HMT Limited, Hyderabad, India (Received June 17, 1991)

Industrial Summary Processing of costly and special materials developed for specific applications needs very careful processing for producing net-shape or near-net-shape components without defects and with controlled microstructures on a repeatable basis. Hence, deformation processing requires sciencebased methodologies to understand and control the workability and microstructural changes under actual processing conditions. In the present paper, atomistic and systems approaches have been used to analyse the deformation processing of RR58 aluminium alloy in the complete range of strain rates and temperatures involved in industrial metal-forming processes. Both of these approaches have yielded comparable results and can be applied quite easily to the processing of different materials.

I. Introduction

The main objectives of the deformation processing of materials are to effect external shape changes and to alter the service properties of the product by controlling the microstructure during deformation. The growing trend of netshape or near-net-shape forming of new and expensive materials necessitates the need to understand various aspects of deformation more thoroughly so that microstructural control can be achieved on a repeatable basis. Several advances have taken place recently to evaluate the deformation processing of materials. In a generalized deformation-processing system as defined by Backofen [ 1 ], the success of mechanical processing depends critically on the microstructure of the initial material, on the geometry of the deformation zone, on the temperature and strain rate employed and on the frictional conditions. From this 0924-0136/92/$05.00 © 1992 Elsevier Science Publishers B.V. All rights reserved.

92 view-point, the atomistic approach, proposed by Frost and Ashby [2] and by Gandhi and Ashby [3] and later extended by Rishi Raj [4] for the development of processing maps defining a "safe" window in the strain rate and temperature field, is of great significance. The third dimension of such a map is the microstructure of the material being processed, which decides the extent of microstructural damage likely to be caused by mechanisms such as cavity formation at hard particles in the low-temperature-high-strain-rate regime, wedge-type microcracks at the grain boundary triple junctions in the hightemperature-low-strain-rate regime, and flow localization due to adiabatic heating at very high strain rates. The properties of the product will be affected greatly by these damaging mechanisms inasmuch as the microstructure of the product will retain the damage caused in the processing stage. The transverse tensile ductility, for example, is dependent sensitively on microstructural damage in the product [ 5 ]. Even though workability is tl{e key aspect of the material during processing, it is complex to define and is dependent on the state-of-stress during forming. Prasad et al. [6] used a systems approach for defining workability in terms of a parameter which is an index of the intrinsic material behaviour and separated this from the state-of-stress related workability. In this approach, also called dynamic materials modelling, the intrinsic ability of the material to dissipate power is the most important criterion to be used in optimizing the deformation process. It considers, essentially, the manner in which the energy is converted at any instance into a form, usually thermal or microstructural, which is not recoverable by the system. The two modes of dissipation are in the form of viscoplastic heat and metallurgical processes and are called Content (G) and Co-content (J) respectively. Prasad et al. [6] have deduced that the partitioning of power between G and J is dependent upon the strain-rate sensitivity (m) of the material. For a given set of processing conditions, they have established the relationship between the Co-content and the strain-rate sensitivity index as follows: J-

a~rn m+l

(1)

where a is the flow stress and t is the strain rate. Further the efficiency of power dissipation (~/) through metallurgical processes was shown to be: 2m

~/=m+l

(2)

The power Co-content and the efficiency of power dissipation serve as the most useful indices for characterizing dynamic material behaviour. Higher efficiency means that a large amount of power dissipates through dynamic me-

93 tallurgical processes leading to return of steady state conditions and high intrinsic workability of the material. Several dynamic metallurgical processes contribute to power dissipation during the hot working of materials and each of these processes has its own characteristic efficiency of dissipation. Often in materials having complex microstructures, two or more mechanisms occur simultaneously. Fracture processes also represent the most efficient means of dissipating power, since very little energy is dissipated as viscous heat. However, such regimes should be strictly avoided in processing. In this paper the deformation processing of RR58 aluminium alloy is studied in the light of the two approaches mentioned above, to provide experimental support of the concepts. 2. Experimental

The RR58 aluminium alloy was prepared using master alloys, rods of 40 m m diameter and 150 m m length being cast in metallic moulds. The cast rods were homogenized at 800 K for 6 hours and were forged at 750 K using a pneumatic hammer to a final size of 25 m m diameter. These forged bars were annealed at 550 K for 6 hours and were then furnace cooled. The alloy has the following chemical composition (mass To): Mg

Fe

Cu

Ni

Si

Zn

Mn

A1

1.23

1.22

2.40

1.31

0.13

0.23

0.02

balance

Cylindrical specimens with an aspect ratio of 1.5 were machined from the annealed bar stock. Isothermal upsetting (compression) testswere carried out at different mean strain rates (0.02, 0.2, 8 and 200 s -I ) in the temperature range of 300 to 800 K utilizingdifferenttest equipment up to a true strain (e) of about 0.7. The room temperature (300 K) mechanical properties (tensile) were obtained using miniature transverse tensile specimens machined from the forged product resulting from the upsetting tests. More details regarding the material, the test equipment, the instrumentation and the experimental procedure are given elsewhere [7,8 ]. 3. Results and discussion

The important information required for studying the deformation processing of the alloy using the atomistic approach and the systems approach is the flow stressat differenttemperatures and strain rates,and the microstructural details of the initialmaterial, in a quantified manner.

94

3.1. Atomistic approach The microstructure of the initial material for upsetting tests revealed fine irregular particles and needle-like particles in the matrix as well as rounded particles at the grain boundaries. The presence of matrix particles was due mainly to precipitation during the slow furnace cooling of the initial material during its preparation. The approximate values of the size and the area fraction of the particles at the grain boundaries are 3 pm and 0.04 respectively, while the values of the size and the volume fraction of the particles in the grain boundary interior (matrix) are 0.4 pm and 0.02 respectively. The average grain size is estimated to be about 25 ~tm. The influence of these particles on the T/Tin 0.4 10

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Fig. 1. Processing map for RR58 aluminium alloy indicating the boundaries for various damage mechanisms. Boundaries A1 and B2 are for pure aluminium (Rishi Raj [4] ) obtained based on the nucleation thresholds for damage mechanisms, while the boundaries D 2 and W 2 are deduced for R R 5 8 alloy on the basis of the ductilityof the product.

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damage mechanisms can be evaluated using the nucleation threshold levelsof strain rate at different temperatures, as demonstrated by Rishi Raj [4] for pure aluminium, assuming the presence of some particlesat the grain boundaries and in the matrix. In the present work, the transverse ductilityof the upset forged product is used as a measure of the microstructural damage that occurred, and is retained, while processing the alloy.The room-temperature ductilityvalues are shown in Fig. I as a function of the processing conditions of strain rate and temperature. The somewhat sudden transition from low ductility to high ductility with increase in the processing temperature indicates clearlythe elimination,

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F i g . 3. Flow stress-true strain curves obtained at a mean strain rate of 8 s - ~ at various test temperatures.

or at least a significant reduction, of the cavitation damage in the matrix at higher temperatures. The transition points are joined to give a matrix particle cavitation damage boundary (marked as D2 ) for this alloy. The transition line obtained by Rishi Raj [4 ] for a similar microstructure in aluminium, based on purely theoretical considerations, is shown also {marked as A1) in the figure. The transition lines are quite comparable, considering the uncertainties that would occur in the quantitative evaluation of microstructures, the considerable deviation of the shape and the size of individual particles in the alloy and the dissolution of some particles at higher temperatures. The observed variation in mechanical properties of the upset forged product

97

has been analyzed in detail elsewhere [8]. The drop in ductility of the product processed at higher temperatures was found to be mainly due to solid solution strengthening in this alloy and should not be mistaken for loss of ductility due to wedge cracking at the grain boundaries while upsetting. The wedge-cracking boundary of the processing map can be obtained from the fracture-surface studies conducted on the tensile samples produced from the upset-forged product [8]. Based on these studies, a line is drawn (marked as W2 ) which delineates the safe-unsafe regions of processing of this alloy at high temperatures. Also shown is the boundary (marked as B2) estimated theoretically by Rishi Raj [4 ] for a similar microstructure in pure aluminium. The large deviation may be due to the limited strain experienced by the RR58 alloy, which may not be sufficient to cause considerable wedge-cracking damage in the present case. Other important damage mechanisms are localized flow due to adiabatic heating and severe localized dynamic recrystallization (in some instances), which are not relevant in the present case of limited strains involved in the upsetting tests.

3.2. Systems approach The flow curves obtained at different test temperatures for mean strain rates of 0.2 and 8 s- 1 are shown in Figs. 2 and 3. The flow curves show typical characteristics of strain hardening and dynamic recovery at lower temperatures, l.O0

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Fig. 5. Variation of flow stress (corrected for the adiabatic temperature rise) with inverse absolute temperature for different mean strain rates, indicating bimodal behaviour of deformation. The bifurcation is indicated by the broken lines.

and dynamic recovery and recrystallization at higher temperatures, the transition from the former to the latter occurring at higher temperature for higher strain rates, as in the case of several conventional metals and alloys. While using the systems approach for the calculation of the process efficiency, one should use as accurate a flow stress of the material as possible. In the present case, the flow stress obtained at an intermediate strain of 0.4 was used, since the influence of frictionon the flow stress at this strain levelis negligible [9,10 ] for the geometry of the specimens used in the present study and also because the uncertainty of elastic compliance at lower strains does not influence the results significantly.The values of the flow stress at this strain level,corrected for the measured adiabatic temperature rise, are shown in Figs. 4 and 5. The influence of temperature and strain rate is quite systematic. However, the data indicate a bifurcation (shown by broken lines) between the low- and the hightemperature ranges. This bimodal behaviour m a y be due to a change in defor-

99

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~. ~o kLI

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Fig. 6. Efficiencyof powerdissipation as a functionof strain rate and temperature. mation mechanism in a certain range of temperature in which it is probability dependent. The systems approach requires the computation of the strain-rate sensitivity (m) as a function of strain rate (~) and temperature. Even though the flow stress data can be represented using much simpler relationships [11,12], it is more appropriate to use the following form of equation: log a=A1 +A2 (log ~) +A3 (log ~)2-bA4 (log ~)8

(3)

Since the flow stress values are available at four different strain rates for each temperature, this third-degree polynomial gives an exact fit at the four strain rates: constants A1, A2, A3 and A4 are thus evaluated. The strain-rate sensitivity (m) is obtained by differentiating eqn. (3) with respect to (log ~): d(log a) m = d(log ~) =A2 +2,43 (log t) + 3 A4 (log t)2

(4)

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Fig. 7. Constant-efficiencycontour map, indicating the bifurcation zone and the optimum processing range for RR58 aluminium alloy. The efficiency values for different processing conditions can now be calculated using eqn. (2), the resulting values being shown in Fig. 6. A more convenient way of expressing t h e m is to represent t h e m as an efficiency contour map, shown in Fig. 7. The map clearly indicates that there is a transition from a low- to a high-efficiency range at about 500 K, which coincides also with the bifurcation range observed (also marked for reference) in the flow stress data presented in Fig. 5: hence in this bifurcation range the microstructure obtainable is probability related. A very high efficiency range can be identified at the lower-right corner of the map - low strain rate and high temperature - which coincides with the wedge-cracking regime shown in Fig. 1. Since fracture mechanisms can have a high process efficiency it is not surprising to observe high efficiency in the case of grain-boundary sliding, which latter causes wedge cracking without producing much viscous heat. Also, the high efficiency range at the top right corner of the map - high strain rate and high temperature could indicate the onset of localized flow due to adiabatic heating, as was discussed by Rishi Raj [4]. The two processing maps obtained using different approaches for the RR58 alloy can be used to eliminate the damage regimes for processing, and one can obtain the regime of high intrinsic workability for this alloy.

101 4. Conclusions (1) Processing maps can be developed successfully using either the atomistic approach or the systems approach, to yield realistic deformation limits: the results of using the different approaches are quite comparable. (2) The approach of developing the processing maps using entirely microstructure-mechanical-property relationships could be very time-consuming and expensive, and could be influenced also by other metallurgical changes that may not be related to deformation. Nevertheless, it is quite a useful check before the processing conditions for an alloy are fixed. (3) The systems approach is quite simple for developing processing maps that can be used for net-shape forming. However, the high-efficiency regimes do not necessarily represent the best processing conditions and need a careful consideration of possible fracture mechanisms which could also be energyefficient. (4) The processing maps obtained by both of the approaches indicate that the RR58 alloy may be processed in the temperature range of 600-750 K and in the strain rate range of 0.2-20 s-1. Acknowledgement The authors would like to acknowledge gratefully the help received from Prof. Y.V.R.K. Prasad, D e p a r t m e n t of Metallurgy, Indian Institute of Science, Bangalore, India.

References 1 2 3 4 5 6 7 8 9 10 11 12

W.A.Backofen, Metall. Trans., 4 (1973) 2679. H.J. Frost and M.F. Ashby, Deformation Mechanism Maps, Pergamon, Oxford, 1982. C. Gandhi and M.F. Ashby, Acta MetaU., 27 (1979) 1565. Rishi Raj, MetaU. Trans., 12A (1981) 1089. H.C.Rogers and L.F. Coffin Jr., Int. J. Mech. Sci., 13 (1971) 141. Y.V.R.K. Prasad, H.L. Gegal, S.M. Doraivelu, J.C. Malas, J.T. Morgan, K.A. Lark and D.R. Barker, MetaU. Trans., 15A (1984) 1883. K.P. Rao and ¥.V.R.K. Prasad, Aluminium, 60 (1984) 184. K.P. Rao and Y.V.R.K. Prasad, Aluminium, 60 (1984) 289. J.A. Schey, T.R. Venner and S.L. Takomana, J. Mech. Work. Technol., 6 (1982) 23. D.L.Baragar, J. Mech. Work. Technol., 14 (1987) 295. K.P. Rao, S.M. Doraivelu and V. Gopinathan, J. Mech. Work. Technol., 6 (1982) 63. K.P. Rao, Unpublished Research, 1991.