INTRODUCTION

Due to environmental reasons, there is an in-creasing demand for Pb-free solders. In fact, the con-straints imposed by politics in Europe make it nec-essary to replace the Pb-containing solders beforethe year 2008.1 Numerous studies on Pb-free soldershave been performed in recent years, most focusingon solder alloys with a melting range below 200°Cthat can replace the Sn-37wt.%Pb eutectic solder.The Sn-based alloys (e.g., Sn-8%Sb) featuring a com-bination of properties that correspond closely tothose in eutectic Pb-Sn have been suggested as adrop-in replacement.2–5 However, for die attach inpower devices, Pb-based solders with a meltingrange above 250°C are used. In this temperature

range, a suitable Pb-free alternative has not yetbeen found. Attempts to adjust the melting range(increase the liquidus and solidus temperatures) ofSn-based solders by appropriate alloying have re-mained unsuccessful until present; the binary phasediagrams of Sn and further metallic elements indi-cate that generally the melting point of pure Sn(232°C) will be lowered by adding elements that arepopular candidates for use in solder alloys. Recently,a Zn-Al-Mg-Ga alloy was introduced as a potentialsubstitute for Pb-containing solders.6 The measuredliquidus and solidus temperatures of the suggestedalloy (Zn-4%Al-3%Mg-3.2%Ga) are given as 347°Cand 309°C, respectively. Some of the physical prop-erties of this alloy compare favorably with those ofPb-5%Sn. Also, the thermal conductivity in the Zn-based alloy is double that of Pb-Sn, and the thermalexpansion rate is lower than in Pb-Sn. However, is-

Journal of ELECTRONIC MATERIALS, Vol. 31, No. 4, 2002 Regular Issue Paper

Zn-Al Based Alloys as Pb-Free Solders for Die Attach

M. RETTENMAYR,1 P. LAMBRACHT,1 B. KEMPF,2 and C. TSCHUDIN3

1.—Darmstadt University of Technology, Materials Science Institute, Petersenstr. 23, D-64287Darmstadt, Germany. 2.—OMG AG & Co. KG, dmc2 Division, D-63403 Hanau-Wolfgang, Germany.3.—ESEC, Hinterbergstr. 32, CH-6330 Cham, Switzerland

The present study examines the potential of Zn-Al based alloys for replacingPb-containing solders used in the die-attaching process. Four alloys (Zn-6wt.%Al, Zn-6wt.%Al-1wt.%Ga, Zn-3wt.%Al-3wt.%Mg, and Zn-4wt.%Al-3wt.%Ga-3wt.%Mg) were cast and characterized in terms of microstructureand mechanical behavior. Based on the criteria of thermodynamics (phase dia-gram extrapolations) and mechanical behavior (flow curves), it was possible toselect the Zn-Al-Ga alloy as best meeting the requirements concerning meltingrange necessary for die attach and mechanical properties. A wire with 0.76-mm thickness was produced successfully by extrusion. However, the alloy, inthe form of wires, is subject to substantial embrittlement after extruding to adegree that further handling becomes impossible. The reason for the embrit-tlement was investigated using x-ray techniques (x-ray diffraction and x-rayphoto spectroscopy (XPS)) and transmission electron microscopy. The tempera-ture dependence of the Ga solubility in the hexagonal close-packed (hcp) Zn-rich phase in combination with the affinity of Ga to the surface of the Al-richand Zn-rich phases is held responsible for the embrittlement. A remedy for theembrittlement could be a reduction of the solubility at high temperatures inthe Zn-rich phase. This will be attempted by adding small fractions of furtheralloying elements in future work.

Key words: Solder alloy, Pb-free, phase diagram, segregation, liquid metalembrittlement

(Received September 12, 2001; accepted December 6, 2001)

278

sues of wettability, thermal fatigue behavior, andbrittleness of the alloy must be addressed. Resultsin the literature suggest that it is impossible to sup-ply this alloy in the form of wires or ribbons.6 Thus,introducing this alloy in the die-attach process en-tails an expensive development effort for die-attachmachines that handle solder in different forms. Itappears that an application of Zn-4%Al-3%Mg-3.2%Ga in actual products is still not imminent.

In the present work, the microstructure evolutionand the mechanical properties of four alloys basedon eutectic Zn-Al, particularly Zn-6wt.%Al, Zn-6wt.%Al-1wt.%Ga, Zn-3wt.%Al-3wt.%Mg, and Zn-4wt.%Al-3wt.%Ga-3wt.%Mg, are studied with theaim of correlating microstructure and propertiesand evaluating the possibilities for influencing theproperties that hinder the direct use of these mate-rials as substitutes for the Pb-containing alloys.

RESULTS AND DISCUSSION

The binary eutectic alloy Zn-6wt.%Al (B) was cho-sen for reference purposes. Its melting point (eutec-tic temperature) is given in the corresponding phasediagram7 as 381°C. Extensive research has beencarried out concerning the ternary system Zn-Al-Mg(most data in the Al-rich corner),8,9 but only limitedinformation is available about the ternary Zn-Al-Gaand the quaternary Zn-Al-Mg-Ga phase diagrams.

Adding Ga to eutectic Zn-Al will lead to a pro-nounced melting range with liquidus and solidustemperatures below the binary eutectic tempera-ture. An experimental phase diagram of the ternaryZn-Al-Ga system is not reported in the literature.9However, as no ternary compounds are expected toexist in this system, tentative phase diagrams havebeen established by means of thermodynamic calcu-lations.9 For the ternary alloy containing Ga (TGa),i.e., Zn-6wt.%Al-1wt.%Ga, the melting range calcu-lated by extrapolating the binary phase diagramsinto the ternary region using the software packageChemSage10 is between 390°C (liquidus tempera-ture) and 348°C (solidus temperature).

The composition of the ternary alloy containingMg (TMg), i.e., Zn-4wt.%Al-3wt.%Mg is close to theternary eutectic composition. With a precise ternaryeutectic composition, the alloy should solidify at aconstant temperature, the ternary eutectic tempera-ture, which is given as 343°C.9

Adding Ga to the TMg alloy will again lead to a pro-nounced melting range rather than a defined meltingtemperature and further lower the liquidus andsolidus temperatures. The melting range of the qua-ternary (Q) alloy Zn-4wt.%Al-3wt.%Mg-wt.%Ga hasbeen determined by differential scanning calorimetry(DSC)6 as lying between 347°C and 309°C.6 Thermody-namic calculations with ChemSage predict a meltingrange between 347°C and 99°C. In these calculations,the liquid fraction drops to less than 5% below 250°C.It appears unlikely that the calculations are grosslyerroneous; we, thus, assume that the low liquid frac-tion has not been detected in the DSC measurements

Zn-Al Based Alloys as Pb-Free Solders for Die Attach 279

in Ref. 6. Also, a liquid fraction in this order is consid-ered in Ref. 6 to be unimportant for practical applica-tions because it is not expected to lead to detachmentof the die during the production process. In the ab-sence of Mg, a Ga content of 3%, desirable for adjust-ing the liquidus temperature and the mechanicalproperties, stabilizes the liquid phase down to theternary eutectic temperature (22°C). The liquidphase, consisting mainly of Ga, is consumed by Mgdue to the formation of the Mg2Ga5 phase.6 The tem-perature-dependent equilibrium composition of theliquid phase of alloy Q, calculated with ChemSage, isshown in Fig. 1.

Liquidus and solidus temperatures of the four al-loys, calculated with ChemSage, and solidus temper-atures measured by DSC are listed in Table I. Theliquidus temperatures could not be determined byDSC measurements because the onset of nucleationand, thus, of latent heat release occurred at largeundercoolings. Identifying the solidus temperatureson the differential temperature plots generated dur-ing the DSC measurement is not trivial inasmuch asthe fcc phase features a miscibility gap (refer to thebinary phase diagram in Ref. 7). The separation ofthe Zn-rich fcc phase into two phases with fcc struc-ture but different composition is accompanied by arelease of latent heat that cannot easily be dis-cerned from the heat release due to the liquid/solidphase transformation.

The as-cast microstructures of the four alloys areshown in Fig. 2. A primary phase (fcc) nucleates be-fore the onset of the coupled eutectic solidification,indicating either retarded nucleation of the Zn solu-tion phase or a slight deviation of the eutectic con-centration from the one given in Ref. 7. In fact,slightly different values can be found in the litera-ture about the composition of eutectic Al-Zn. In thedata file of the OMG/dmc2 database used for theChemSage calculations, the eutectic concentration isZn-5.3wt.%Al rather than Zn-6wt.%Al as in Ref. 7.

The four aforementioned alloys were cast in in-gots (12-mm diameter and 100-mm length) and ma-

Fig. 1. Equilibrium concentrations in the liquid phase of alloy Q as afunction of temperature, calculated with ChemSage.

280 Rettenmayr, Lambracht, Kempf, and Tschudin

chined to cylinders of 10-mm diameter and 15-mmlength. The cylinders were first deformed in com-pression tests at 310°C at a strain rate of 0.003 s21.The flow curves generated in these tests (Fig. 3) pro-vide a means of judging the deformability for wireextrusion.

Comparison with the flow curve of an alloy that isat the limit of producibility in wire form (Alloy J inFig. 3) shows that the yield strengths of two of thealloys (Q and TGa) are below the critical value. All ofthe alloys feature a strong decrease of true stress attrue strains above 0.1. This can be attributed to oneor both of the following reasons. First, plastic flow isfacilitated by an increase of the effective stress inthe active gliding systems of the sample due to crackformation, and second, drastic changes in the mi-

crostructure due to the deformation may change themechanical properties significantly.

As macroscopic cracks at the surface were not ob-served in the compression test, the reason for thedrop of the flow stress is expected to lie in the mi-crostructure. Metallographic samples were preparedfrom the deformed specimens. The microstructuresof the four alloys after deformation are shown inFig. 4. The microstructures appear to be essentiallycrack free. The length scale of the microstructure isreduced drastically by one to two orders of magni-tude, yielding the observed drop in flow stress.

The flow curves, as shown in Fig. 3, provide a cri-terion for a sufficiently low flow stress for wire pro-duction. However, there is only preliminary informa-tion concerning a sufficiently high ductility in the

Table I. Liquidus and Solidus Temperatures of the Zn-Based Alloys

Designation Composition (wt.%) TL (°C) TS (°C) TL (°C) TS (°C)

ChemSage DSC Measurements

B Zn-6Al 391 381 — 380TGa Zn-6Al-1Ga 390 351 — 366TMg Zn-4Al-3Mg 362 346 — 341Q Zn-4Al-3Ga-3Mg 353 99 3476 3096

Fig. 2. As-cast microstructures of the alloys: (a) Zn-6wt.%Al, (b) Zn-6wt.%Al-1wt.%Ga, (c) Zn-3wt.%Al-3wt.%Mg, and (d) Zn-4wt.%Al-3wt.%Ga-3wt.%Mg.

Zn-Al Based Alloys as Pb-Free Solders for Die Attach 281

materials. Even though the majority of the alloystested that passed the flow curve criterion could beextruded to wires, in some examples, the wires weretoo brittle to be handled or fractured into smallpieces of a few centimeters length during extrusion.

This latter effect was observed in the case of the Qalloy. Thus, the findings of the research group in Ref.6 were confirmed with the methods and equipmentavailable for the present study.

Alloy TGa could be extruded to a wire that waswound automatically around cores with 2-cm diame-ters. Thus, it appears that one of the shortcomings ofalloy Q, particularly the exceedingly high brittle-ness, could be overcome by decreasing the Ga con-tent and avoiding Mg as an alloying element. How-ever, an unexpected embrittlement of the wire wasobserved within 24 h after extrusion to a degree thatrendered the wire unsuitable for use in practical ap-plications because it cannot be handled by an auto-matic soldering device.

Ternary phases that are not present in the bi-nary phase diagrams are not likely to be found inthe Zn-Al-Ga system. The precipitation of furtherphases acting as a cause for the embrittlement isnot expected and has also been excluded using x-ray diffraction analysis and transmission electronmicroscopy diffraction patterns; phases other thanthose described in the calculations could not beidentified. Further investigations were carried outby studying the morphology and the composition ofthe fracture surface. The fracture surface of anewly extruded wire consists of dimples that are

Fig. 3. Flow curves of the four alloys, generated in compressiontests at 310°C, together with the flow curve of reference alloy Sn-25wt.%Ag-10wt.%Sb (Alloy J), measured at T 5 190°C.

Fig. 4. Microstructures of the alloys after deformation in compression tests: (a) Zn-6wt.%Al, (b) Zn-6wt.%Al-1wt.%Ga, (c) Zn-3wt.%Al-3wt.%Mg,and (d) Zn-4wt.%Al-3wt.%Ga-3wt.%Mg.

282 Rettenmayr, Lambracht, Kempf, and Tschudin

characteristic for ductile materials (Fig. 5a),whereas the fracture surface of the 24-h cold-agedwires (Fig. 5b) clearly shows the intergranular na-ture of crack propagation. The XPS analysis of thefracture surfaces shows a significant compositionchange, whereas the surface of the newly extrudedwire is of bulk composition, the 24-h cold-agedwires reveal a strong enrichment of Ga of up to 32wt.% Ga. Thus, the Ga enrichment could be identi-fied as being responsible for the embrittlement ofthe wire.

The Ga is known to cause liquid metal embrittle-ment (LME) in Al by dramatically reducing the co-hesion between Al grains.11 It is also mentioned inthe literature that there is a weak effect of LMEcaused by Ga in Zn.12 The driving force for the pene-tration of liquid Ga into the grain boundaries is as-sumed to be the reduction of the grain boundary en-ergy rather than the chemical potential gradientthat would lead to grain boundary diffusion.13 How-

Fig. 5. Fracture surfaces of Zn-6wt.%Al-1wt.%Ga wires (a) failed after extrusion and (b) 24 h after extrusion.

ever, LME is not likely to be responsible for the em-brittlement of the solder wire in the present case.Equilibrium calculations with ChemSage indicatethat the liquid phase is not stable below 350°C.Under ideal nonequilibrium (Scheil) conditions, asmall fraction of liquid will persist down to very lowtemperatures (room temperature), as shown in Fig.7. However, the predicted liquid fraction is less than1% at temperatures lower than 220°C. Under real-istic conditions, the low liquid fraction predictedwith the Scheil approximation is further reduced byback diffusion.

If the liquid phase (consisting at low tempera-tures mainly of Ga; refer to the solidification path inFig. 8) was responsible for the embrittlement, thiseffect should be present immediately after solidifi-cation without the observed delay of 24 h after wireextruding. Furthermore, the evolution of the mi-crostructure from this point on should be towardequilibrium, and the liquid fraction should disap-

Fig. 6. The Ga concentration in the (nonstoichiometric equilibrium) phases as a function of temperature in the alloys (a) Q and (b) TGa, calcu-lated with ChemSage.

Zn-Al Based Alloys as Pb-Free Solders for Die Attach 283

pear in a time span determined by the diffusion ki-netics.

The solubility of Ga in hcp Zn (approximately 1wt.%7) and fcc Al (20 wt.%) is as high as, or substan-tially higher than, the Ga content of the alloy. How-ever, the solubility of Ga in the different phases inthe ternary or quaternary alloys will be strongly in-fluenced by the composition of the alloys. This isshown in Fig. 6 where the Ga concentrations in theequilibrium phases (alloy TGa) and in the nonstoi-chiometric phases (alloy Q) are plotted as a functionof temperature. As opposed to the preceding LMEexample, in the alloys TGa and Q, the Ga is dis-solved in the alloy after solidification, and overall,the Ga content is below the solubility limit (becausephases other than fcc and hcp are not stable at roomtemperature).

For pure Zn, a certain amount of surface segrega-tion of Ga has been observed and used for separat-

ing radioactive Ga isotopes from Zn targets,14 indi-cating an interface activity of Ga that may also leadto grain boundary segregation. Because surface seg-regation and grain boundary segregation are associ-ated with a reduction of the interfacial energy, andthe average free energy of grain boundaries is gen-erally about 1/3 to 1/2 of that for a free surface, it isclear that grain boundary segregation should be aless pronounced effect than that of surface segrega-tion. In fact, it has been stated that embrittlementdue to segregant elements can be characterizedusing the difference of sublimation enthalpies15 as ameasure. According to this approach, grain bound-ary segregation of Ga is not to be expected in Zn be-cause the difference between the sublimation en-thalpies of the matrix element and the segregantshould be positive, whereas in the present case, it isnegative. Furthermore, the small difference ofatomic radii of Ga and Zn suggests that there is onlya small elastic-strain energy due to the solute solu-tion, which is directly related to grain boundary seg-regation (Ref. 16). Thus, only a small driving forcefor grain boundary segregation is present.*

The preceding reasoning on grain boundary segre-gation is directly applicable only to binary alloys. Inthe present case, the ternary composition of thealloy may not be negligible. In addition to the Znphase, in TGa, as a second phase, the fcc Al phase ispresent. At temperatures above the miscibility gapof the fcc phase (Ref. 7), i.e., just after solidification,the fcc phase contains a high concentration of Zn(above 80 at.%). Below the decomposition tempera-ture (279°C), the amount of Zn in the fcc phase issignificantly reduced to less than 15 at.% below150°C. It can be assumed that the surface activity ofGa on Al, influenced but not significantly reduced bythe presence of Zn, will still play a role.

In Fig. 6b, it can be seen that the equilibrium Gaconcentration in the solid phases is temperature de-pendent, whereas in the Zn phase, the Ga solubilitydecreases with decreasing temperature (as is com-mon for the majority of solute elements in metal al-loys), and in the fcc phase, it increases strongly. Inthe TGa alloy, the decay of the fcc phase into the Ga-poor fcc and Zn phases was observed in DSC meas-urements at some undercooling, above 200°C. Thefact that this reaction occurs at this given tempera-ture indicates that the kinetics is still sufficientlyfast to maintain the sample in a state not too farfrom equilibrium at 250°C. Assuming that the kinet-ics slows down with further decreasing temperatureto a degree that the concentration distribution is notdrastically changed any more in the time scale ofthe cooling process, the Zn phase, i.e., the majorityphase, will be supersaturated with Ga at room tem-

Fig. 7. Calculated (by ChemSage) liquid fraction as a function oftemperature in the four alloys and calculated liquid fraction as afunction of temperature assuming Scheil conditions in the alloy TGa.

Fig. 8. Calculated liquidus projection of the system Zn-Al-Mg withsolidification paths for the alloy TGa under equilibrium (beginningand end points: d) and Scheil conditions (j).

*Using values of the atomic radius of Ga from compila-tions, such as periodic tables, is somewhat problematic.For a more detailed discussion of atomic radii of Ga in Alrefer to D.I. Thomson, V. Heine, M.C. Payne, N. Marzari,and M.W. Finnis, Acta Mater. 48, 3263 (1999).

perature. This implies that a significant fraction ofthe Ga needs to diffuse into the fcc phase to achievethe final equilibrium state. Comparing the diffusioncoefficients of Ga in (pure) Al and (pure) Zn17 atroom temperature reveals that diffusion in Zn isfaster by six orders of magnitude. This means thatGa diffusion out of the Zn phase is still possible,whereas the absorption of the Ga by the fcc phasewill not occur in the same time span.

The embrittlement of the TGa solder wire is, thus,taken to be a combined effect of the reduction of thesupersaturation of Ga in the Zn phase and the sur-face activity of Ga on the fcc Al and hcp Zn phases.The time scale of the embrittlement should then bedictated by the kinetics of diffusion of Ga out of thehcp Zn phase. This is verified by using an analyticalsolution of the diffusion equation.18 Approximatingthe polygonal shape of the grains (compare Fig. 5b)as spheres with a diameter of d 5 1.2 mm, the evolu-tion of the Ga concentration in the Zn-rich phasecan be estimated. The initial bulk concentration isset to the concentration calculated with ChemSageat 250°C (ci 5 0.85 wt.%), and the concentration atthe boundary is set to the value at room tempera-ture (ce 5 0.3 wt.%). It is assumed that the Ga atomsdiffusing out of the Zn phase can be easily accommo-dated in the grain boundaries due to the presence ofthe fcc phase; the boundary concentration is, thus,kept constant at the equilibrium value. Under theseconditions, the evolution of the Ga concentration c inthe Zn phase at a given position r is described by:18

where a is the radius of the sphere (i.e., grain), n isthe index of the series, D is the diffusion coefficient,and t is the time. A graphical representation of con-centration profiles generated using the precedingequation with the diffusion coefficient at room tem-perature is shown in Fig. 9.

Within 24 h, approximately 20% of the excess Gadiffuses out of the Zn phase. With the given radius ofthe grains, an atomic radius of Ga of rGa 5 1.35 Å (cal-culated using A 5 (rNA ? (2rGa)3), A as the atomicmass, r as the density, and NA as the Avogadro num-ber,16 this quantity of Ga results in a monoatomiclayer on the Zn phase. Thus, the time scales for the re-duction of the supersaturation by diffusion and theembrittlement of the wire correspond. At present, itappears to be impossible to quantify the role of grainboundary segregation as a further mechanism favor-ing embrittlement in the Zn-Al-Ga system. Theamount of Ga that remains in the grain boundaryafter long times is defined by the energy balance of Gaatoms in the bulk and the grain boundary. Theamount that will eventually diffuse into the Al-rich

nn ra

Dna

n−

( ) sin expπ π1 2 2

2

c c c c ar

c ci e i e in

= + − + −−

∞

∑( ) ( )π

21

284 Rettenmayr, Lambracht, Kempf, and Tschudin

phase cannot be determined due to the low diffusioncoefficient of Ga in Al. Irrespective of the effects ofgrain boundary segregation, the mechanism for theembrittlement is diffusion of Ga out of the Zn phase.In further work, the possibility of influencing the solu-bility of Ga in Zn and its temperature dependence, aswell as the interfacial energy of the Al phase by intro-ducing further alloying elements will be investigated.

SUMMARY

The Zn-Al-Ga(-Mg) alloys are candidates for Pb-free solders applicable in the temperature rangeabove 250°C. A number of alloys have been evalu-ated that meet the majority of requirements thatare given by the combination of conditions duringthe soldering process and application of the compo-nents. The shortcoming of alloy Q, as described inRef. 6, namely, the too high brittleness for wire pro-duction, could be overcome by reducing the Ga con-tent and avoiding Mg as an alloying element. How-ever, wires of the alloy TGa showed a substantialembrittlement within 24 h after extrusion. Themechanism of the embrittlement is understoodsemiquantitatively. The combination of thermody-namic calculations with a small number of experi-ments has been shown to be a powerful tool for alloydevelopment in the given system.

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Fig. 9. Evolution of the Ga concentration in the Zn phase as a func-tion of time at room temperature.

Zn-Al Based Alloys as Pb-Free Solders for Die Attach 285

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