Enhanced AW3003 Aluminum Alloys for HeatExchangers

M. Pokova,1,2 M. Cieslar,1 J. Lacaze2

1Charles University in Prague, Faculty of Mathematics and Physics, Ke Karlovu 5, 121 16Prague 2, Czech Republic.2Institut National Polytechnique de Toulouse, ENSIACET, 6 allee Emile Monso, 31029 Toulousecedex 4, France.

Abstract. This paper comprises a survey of the microstructure, phase composition,physical properties and processing of Al-Mn based aluminum alloys used for theproduction of automotive heat exchangers. After twin-roll casting the aluminumfoil undergoes during fabrication several annealing and cold-rolling steps. Allthese steps have influence on properties of the final product. During the thermo-mechanical treatment precipitation of dispersoids and coarsening of primary phasesoccur. Dominant phases are α-Al15(Fe,Mn)3Si2 and Al6(Fe,Mn). Deformed grainsare replaced by new ones during recrystallization which results in softening of thematerial. Enhanced properties can be achieved by adding minor alloying elementslike zirconium or chromium. Zirconium can form fine precipitates which retardrecrystallization and stabilize microstructure but their inhomogeneous distributioncan decrease this favourable effect.

Introduction

Thanks to the low density and good strength, formability, thermal conductivity, corrosion resistance,weldability, recyclability etc., aluminum and its alloys belong to the most widely used metals in manybranches of the industry. Aluminum-manganese alloy AW3003 is commonly used as fins in automotiveheat exchangers [Carlberg et al., 2006]. Its standard composition is given in Tab. 1. Recent aim is toreduce the thickness of the foils from which the final exchanger is formed from 100 µm to 50 µm. Thiswill contribute to the reduction of mass and increased efficiency. However, to achieve such a reductionof thickness, currently used material has to be modified: either by changing the chemical composition orthe processing route.

Table 1. Composition of commercial AW3003 alloy used in automotive heat exchangers.AW 3003 Mn Fe Si Cu Zn[wt.%] 1.0-1.5 ≤0.7 ≤0.6 0.05-0.2 ≤0.2

Twin-roll casting

The foils for heat exchangers can be produced either by direct-chill casting (DC) or twin-roll casting(TRC). The use of TRC instead of DC leads to a lower energy and work consumption, higher productionrate and thus lower foil costs, but also to the different initial microstructure. Foils produced by TRCexhibit higher solid solution supersaturation, finer dispersion of primary phases and finer grains elongatedin the casting direction. Therefore, production of TRC materials requires modifications of the processingroute which is used for DC materials.

During twin-roll casting the molten metal is directed through a ceramic tip between two water cooledlubricated rolls. By combination of the solidification and rolling process strips 10–0.5 mm in thicknessare produced [Yun et al., 2000]. The solidification rate around 100 K/s is higher at the surface than inthe centerline of the foil [Sun et al., 2006]. Higher casting rate increases the productivity but can causea macro-segregation in the central part of the foil (Fig. 1). The two limiting surfaces of the TRC foilsolidify first inducing the solute enrichment of the liquid. Solute-rich areas with eutectic features parallelto the casting direction are formed when the two solidification fronts reach the foil center. According to[Birol, 2009d] the segregation is more pronounced in materials with a higher content of alloying elementsand longer freezing range and segregates cannot be removed by a homogenization treatment.

The microstructure of the TRC foils varies through the thickness of the foil: near the surfacethe number density of primary particles is higher and grains are smaller as compared to the center of

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POKOVA ET AL.: ENHANCED AW3003 ALUMINUM ALLOYS FOR HEAT EXCHANGERS

Figure 1. The eutectic centerline segregation in the TRC foil caused by a high casting rate when theliquid enriched in solutes solidifies enclosed by two already solidified skins.

Figure 2. Light optical micrograph of the grain structure of a twin-roll cast AW3003 alloy. Surfacegrains are deformed and elongated in the casting direction (left), grains in the center are coarser (right).

the foil [Birol, 2009a]. Number of dispersoids precipitated during the heat treatment is higher in thecenter [Karlık et al., 2009]. The central grains are more or less equiaxed whereas the surface grainscan be elongated in the casting direction (Fig. 2) or very fine thanks to the recovery and dynamicrecrystallization caused by hot-rolling component of twin-roll casting [Birol, 2009a,c].

Manufacturing of heat exchangers

After twin-roll casting the aluminum foils undergo several annealing and cold-rolling steps. In the as-cast state the homogenization at temperatures 550–600 ◦C can be applied to reduce the microsegregationand relax the strain introduced into the matrix by the casting process. As a consequence of homoge-nization, material softens more easily and at lower temperatures as compared to the inhomogenized one[Birol, 2009b]. Precipitation annealing can lead to the formation of fine dispersoids and transformationof morphology and/or size and distribution of the primary particles and is usually applied at temperatureranging between 300 and 500 ◦C.

Cold-rolling decreases the foil thickness, increases the dislocation density and the stored deformationenergy in the material and causes grain elongation in the rolling direction. From the as-cast thicknessaround 10 mm to the final thickness 0.1 mm foil undergoes several rolling steps. At the intermediatethickness around 1 mm the recrystallization annealing is applied to release the stored energy and softenthe material to enable further rolling [Pokova, 2010].

As a final step, the fin is formed from the foil cold-rolled to the final thickness and is brazed tothe cladding from an aluminum-silicon alloy of the type AW4343. Brazing is carried out at temperaturenear the melting point of the alloys and can evoke grain coarsening, deterioration of strength, saggingand inhomogeneous structure: thanks to the diffusion of silicon from the cladding to the core material aband of dense precipitates can be formed at the interface [Lacaze et al., 2005; Dehmas et al., 2006].

Final mechanical properties depend on the concentration, size and distribution of both primary and

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secondary particles and are thus influenced by the temperature range and holding time of all annealingsteps. To acquire desirable strength and formability, all processing steps require thorough understanding.

Precipitation and phase transformation

The initial microstructure of TRC aluminum alloys AW3003 is influenced by the content of alloyingelements and parameters of casting (e.g. the foil thickness, casting rate, roll diameter, load). Generallyit consists of the primary particles containing aluminum, manganese, iron and/or silicon at dendriteboundaries [Cieslar et al., 2007] and supersaturated matrix. Thanks to a high cooling rate, the matrixis after solidification in a non-equilibrium state.

The primary particles in aluminum alloys formed by eutectic reactions are Si, cubic α-phase with thenominal composition Al12−15(Fe,Mn)3Si1−2 and orthorhombic Al6(Fe,Mn) phase [Cieslar et al., 2008].Which phase is dominant depends mainly on silicon content and solidification rate [Birol, 2009a]. Iron andmanganese can substitute for each other in the later two phases [Shabestari, 2004]. Crystal structureof α-phase depends on the Fe:Mn ratio: for higher Mn content the structure is simple cubic, withincreasing iron content it transforms into body-centered cubic [Li and Arnberg, 2003b]. During annealingAl6(Fe,Mn) transforms into α-phase thanks to the diffusion of silicon from the solid solution. Decreaseof Fe:Mn ratio during the transformation implies that diffusion of Mn is of importance too. α-phasenucleates at the Al6(Fe,Mn)-matrix interface and propagates into Al6(Fe,Mn) as an eutectoid reactionfront [Alexander and Greer, 2002]. According to [Dehmas et al., 2005] the transformation in DC materialsstarts around 350 ◦C, according to [Li and Arnberg, 2003a] around 400 ◦C. Transformation is controlledby the nucleation, starts rapidly and follows simple relaxation kinetics [Alexander and Greer, 2002].

In the course of the heat treatment, eutectic colonies of primary particles break up and spheroidize;at temperatures above 560 ◦C their size keeps increasing but their number density decreases [Li andArnberg, 2003a].

During annealing small dispersoids precipitate from the supersaturated aluminum matrix (Fig. 4).They are identified as α-phase and have lower Fe:Mn ratio than the primary particles [Li and Arnberg,2003a; Dehmas et al., 2005]. Precipitation in non-heat treatable Al-Mn-based alloys occurs mainlyheterogeneously on dislocations, subgrain and grain boundaries [Cieslar et al., 2008].

Temperature range of the most intensive precipitation depends on alloy composition, thermo-mechanical history and heating rate during annealing. Precipitation can be suppressed by heating ratesas high as 20 K/s [Dehmas et al., 2002]. Precipitation is usually observed between 300 and 500 ◦C[Pokova et al., 2009; Cieslar et al., 2007; Dehmas et al., 2002]. Size of dispersoids increases with increas-ing temperature and holding time. At temperatures above 540 ◦C [Pokova et al., 2009] the smaller onesstart to dissolve back into aluminum matrix thanks to the increased solubility of manganese. Mn fromdissolved dispersoids can then contribute to the coarsening of primary particles [Li and Arnberg, 2003b].

If a two-step annealing is applied, final dispersoid distribution is influenced by both annealingtemperatures. If the second temperature is lower than the first one, dispersoids formed during thesecond step nucleate heterogeneously on coarser particles which precipitated during the first step andtheir number density is low. In the oposite case fine dense dispersoids formed at lower temperatureredissolve and spheroidize at high temperature [Huang and Ou, 2009].

The precipitation kinetics can be easily monitored by conductivity or resistivity measurements. Onresistivity annealing spectra temperature areas of main precipitation and depletion of supersaturatedsolid solution are represented by peaks in positive values whereas dissolution of constituents back to thealuminum matrix corresponds to valleys (Fig. 3). The contribution to resistivity is much higher fromsolute atoms dissolved in matrix than from other crystal defects; manganese atoms in matrix affectsresistivity more strongly than iron and silicon [Cieslar et al., 2007; Li and Arnberg, 2003a]. According to[Li and Arnberg, 2003a] variations of manganese in the solid solution during heating are caused mainly bythe formation and dissolution of dispersoids and are not influenced significantly by the transformations ofthe primary particles: at low temperatures diffusion rate of Mn is low to diffuse to the primary particles;on the other hand, at higher temperatures manganese which enriches the solid solution (causing resistivityincrease) comes from fine dispersoids which can be more easily dissolved than the primary phases.

Recrystallization

If a metal is deformed by a hot or cold working, dislocations are generated and stored in thematerial (Fig. 4). Dislocations are non-equilibrium lattice defects and thus the material tends to removethem in order to achieve thermodynamic equilibrium [Gottstein and Shvindlerman, 2010]. If we applya heat treatment, recovery and recrystallization occur, leading to softening of the material. During

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-800

-600

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0

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-1/

ρ 0 d

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With ZrNo Zr

Figure 3. Left: Resistivity change during isochronal annealing, peaks correspond to precipitation anddepletion of matrix of solutes, valley represents dissolution of constituents back to the solid solution.The spectrum is different for as-cast and cold-rolled material. Right: Vickers hardness test: Shift ofrecrystallization temperature during isochronal annealing by 50 ◦C in material with Al3Zr dispersoids.Recrystallization is connected with profound decrease of HV.

Figure 4. TEM micrographs of a cold-rolled material AW3003 with high density of dislocations (left)and fine dispersoids formed during annealing to 540 ◦C distributed between coarser primary particles ineutectic colonies (right).

recovery dislocations annihilate and rearrange themselves into dislocation cells. Recrystallization resultsin formation of new dislocation-free grains and their subsequent growth (Fig. 5).

Temperature and extent of the recovery and recrystallization and final grain size are influenced bythe thermo-mechanical history of the material. Large particles in the matrix promote recrystallization byproviding additional nucleation sites—particle simulated nucleation [Engler, 1998]. On the other hand,finely dispersed precipitates of suitable size and spacing can retard recrystallization by stabilizing thedislocation structure and pinning moving dislocations (Eq. 1). As a result of pinning critical radius ofrecrystallization nuclei increases and a coarse grained material is obtained [Birol, 2009c].

The higher is strain induced by cold-rolling the higher is the stored deformation energy and thusdriving force for recrystallization. However, if the strain exceeds certain level, recovery becomes dominantprocess, the driving force for recrystallization is reduced and the recrystallization resistance is increased[Birol, 2008].

Recrystallization can be easily monitored by microhardness measurements (Fig. 3) and light mi-croscopy observations [Slamova et al., 2004].

Addition of zirconium

By adding small amounts of zirconium into aluminum alloys metastable coherent dispersoids Al3Zrwith cubic L12 structure can be formed (Fig. 6). They stabilize the microstructure and pin moving grainboundaries by exerting Zenner drag (Eq. 1) and thus retard recrystallization (Fig. 3). The Zener drag Zis proportional to the grain boundary energy γGB (∼ 0.3 J/m2 for high angle boundaries in aluminum)and increases with the increasing volume fraction f and decreasing radius r of the dispersoids.

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Figure 5. Cold-rolled material with deformed grains elongated in the rolling direction and materialannealed at 440 ◦C with new recrystallized grains, the rolling direction still evident.

Figure 6. TEM micrographs of metastable cubic Al3Zr dispersoids with average size 10 nm, dark field(left) and inhomogeneously distributed Al3Zr: high density in the left upper corner, no dispersoids inthe right low corner, bright field (right).

Z =3fγGB

2r(1)

The highest recrystallization resistance in direct-chill cast material was reported by [Jia et al., 2008]after two-step annealing at 250 and 450 ◦C with slow heating rate. However, Al3Zr dispersoids areinhomogeneously distributed in the aluminum matrix (Fig. 6) with periodicity corresponding to thesecondary dendrites arm spacing [Jia et al., 2007]. This inhomogeneity is caused by microsegregationof Zr during casting. Areas with low density of the dispersoids are still prone to recrystallization.Homogenization before precipitation annealing only decreased the driving force for the precipitation andthe final volume fraction of dispersoids so that homogenized material recrystallized at lower temperaturethan the inhomogenized one [Jia et al., 2007].

Conclusion

Final mechanical properties of a heat exchanger are influenced by all thermo-mechanical steps under-taken during its manufacturing. To fulfil industrial requirements, all steps have to be carefully monitoredand adjusted. During brazing at high temperatures, when the grain coarsening and decrease in strengthcan occur, recrystallization resistance is required. That can be achieved by dispersion of fine particleslike Al3Zr. The kinetics of their precipitation depends on the annealing temperature, heating rate, com-position, solute content available in the matrix, primary particles distribution and lattice defects inducedby the cold-rolling. However, their pinning effect is reduced by coarser particles which serve as nucleationcenters for recrystallization and whose formation during heat treatment is inevitable.

Acknowledgments. The work was financially supported by grants GAUK 106-10/251153 and SVV-2010-261303 and scholarships from French government and AL INVEST Bridlicna, a.s., Czech Republic, whoalso provided all experimental materials.

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