Aluminum Alloy 5052 and L-C Steel by Laser Roll Welding

8/13/2019 Aluminum Alloy 5052 and L-C Steel by Laser Roll Welding

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ABSTRACT. Low-carbon steel and alu-minum Alloy 5052 sheets were diffusion

welded using different conditions of time,temperature, and pressure in order to un-derstand the kinetics of diffusion and themicrostructure of the interface layer. Highpressures accelerated intermetallic com-pound formation at the interface. The in-terface layer consisted of aluminum-richbrittle intermetallic compounds (FeAl3

and Fe2 Al

5), which made the joints brittle.

In laser roll welding, steel and alu-minum sheets in a lap-joint configurationare subjected to laser heating and imme-diate rolling for intimate contact. Laserheating provides high temperature in ashort time for melting of the aluminumalloy and diffusion. It results in formationof a thin interface layer. When the tem-perature is above 1473 K (1200°C), for-mation of iron-rich (Fe-rich) intermetalliccompounds (FeAl and Fe3 Al) is encour-aged. Travel speed and roll pressure were

varied and resultant joints were character-ized by optical microscopy, electron-

probe microanalysis (EPMA), scanningelectron microscopy (SEM), X-ray dif-fraction, and tensile shear testing. Thelaser roll welded interface layer containsbrittle aluminum-rich (Al-rich) inter-metallic compounds on the aluminum sideand slightly ductile Fe-rich intermetalliccompounds on the steel side. As the travelspeed increases, thinner interface layersare formed and the percentage of Fe-richintermetallic compounds in them in-creases. This increases the shear strengthof the joints from 11.0 to 55.9 MPa. Inter-face layers with a thickness of 4 to 5 µm,

containing 25 to 40% Fe-rich intermetal-lic compounds, gave maximum shearstrength.

Introduction

Hybrid structures of aluminum alloyand steel are suggested for reducing the

weight of automobiles to improve fuel ef-

ficiency and control air pollution (Refs. 1,2). Therefore, joining steel and aluminumalloy in different shapes is receiving atten-tion. However, iron and aluminum are notcompatible metals as far as fusion weldingis concerned. The reason for this is attrib-uted to the large difference between theirmelting points (933 K for Al and 1811 K for Fe), the nearly zero solid solubility of iron in aluminum, and the formation of brittle intermetallic compounds such asFe2 Al5 and FeAl3. Further, differences intheir thermal properties — like expansioncoefficients, conductivities, and specificheats — lead to internal stresses after fu-sion welding. Therefore, fusion welds of iron and aluminum suffer from heavycracking with brittle failure in service.

The Fe-Al phase diagram is shown inFig. 1 (Ref. 3). The intermetallic com-pounds present on it are grouped as Fe-rich compounds (FeAl and Fe3 Al) and Al-rich compounds (FeAl2, Fe2 Al5, andFeAl3). Table 1 summarizes their compo-sition and crystal structures (Refs. 3, 4). In

addition to these stable compounds,metastable compounds (FeAl6, Fe2 Al9,and FeAl x ) are also shown in Table 1 (Ref.5). Mechanical properties like hardness,fracture toughness K IC values by indentionfracture method, compressive strength,and strain under compression of cast in-termetallic compounds are shown in Table2 (Ref. 6). While Al-rich intermetalliccompounds are hard and brittle (0% com-pressive strain), Fe-rich intermetalliccompounds show slight ductility (0.45 and0.8% compressive strain). Iron-rich inter-metallic compounds also possess higher

compressive strength than Al-rich inter-metallic compounds.

Diffusion Welding

The effect of temperature on diffusioncoefficients is shown in Fig. 2 (Ref. 7). Dif-fusion of iron into aluminum is muchfaster than vice versa. The condition is fa-

vorable for forming brittle Al-rich inter-metallic compounds in the interface layerbecause iron atoms are fewer in numberafter migrating into the aluminum side. As

Al-rich intermetallic compounds are brit-tle, diffusion-welded joints show poorstrength unless insert material is used tocontrol their formation (Refs. 8, 9).

Kinetics of Intermetallic Compound Formation

When a dissimilar metal combinationinvolves formation of intermetallic com-pound at the interface, it takes place intwo stages. Initially, supersaturated solidsolution is formed due to migration of atoms across the interface. When compo-sition of supersaturated solid solutionreaches to a sufficient level, it transforms

into an intermetallic compound (Ref. 10).Therefore, the time before the inter-metallic compound appears at the inter-face is called “incubation time” during

which supersaturated solid solution is en-riched by diffusing atoms. Activation en-ergy for the formation of supersaturatedsolid solution during the incubation timeis denoted as ESol. Effective activation en-ergy for the subsequent transformation of supersaturated solid solution into inter-metallic compound is denoted as E Eff . Thetotal activation energy for the formationof intermetallic compound ( E IMC) can be

defined as

E IMC = ESol + E Eff (1)

This means that E IMC is greater than ESol.When there is a large difference in dif-

fusion rates (as in the case of Fe-Al), thesolute atoms and solvent lattice will be de-cided and the formation of supersaturatedsolid solution will be favored on one sideof the interface. The rate of formation of supersaturated solid solution ( Ñ Sol) withinthe incubation time will simply depend onthe diffusion of solute atoms in the solvent

lattice. The Arrhenius equation for the

Joining of Aluminum Alloy 5052 andLow-Carbon Steel by Laser Roll Welding

A combination of laser heating and roll welding is suggested to join aluminum Alloy 5052 and low-carbon steel sheets

BY M. J. RATHOD and M. KUTSUNA

M. J. RATHOD and M. KATSUNA are with De- partment of Material Processing Engineering ,School of Engineering, Nagoya University,

Nagoya, Japan.

KEY WORDS

Dissimilar MetalsDiffusion WeldingKirkendall PorosityLaser Roll WeldingLap JointIntermetallic Compound



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formation of supersaturated solid solution will be

Ñ Sol = kSol exp[–ESol /RT] (2)

where kSol is rate constant in incubationtime prior to intermetallic compound for-mation, R is gas constant, and temperatureof diffusion is T in K.

Similarly, rate of transforming the su-

persaturated solid solution into inter-metallic compound ( Ñ Eff ) is given by

Ñ Eff = k Eff exp [–EEff /RT] (3)

where, k Eff is rate constant for convertingsupersaturated solid solution into inter-metallic compound. Effective activationenergy, E Eff , depends on the factors in-

volved in this transformation.

E Eff = nW + Ecr + ne (4)

where nW is the activation energy essential

for localized fluctuations in composition, Ecr is the activation energy involved in re-structuring of the crystal lattice when theintermetallic phase is formed, and ne is thedeformation energy of the matrix sur-rounding the nucleus.

Because of the difference in factors involved in formation of supersaturatedsolid solution (involving atomic diffusion)and transformation of supersaturatedsolid solution into intermetallic compound(Equation 4), the rate constants kSol and k Eff

are different.The value of activation energy for the

formation of supersaturated solid solution( ESol) will always remain constant (in caseof Al diffusing in Fe, it is 135 kJ/mol). Oncethe transformation of supersaturated solidsolution into intermetallic compoundstarts, its rate will vary depending on thediffusion of solute atoms into the inter-metallic compound layer formed. In this

way, in addition to the previously statedfactors (Equation 4), another factor (dif-fusion coefficient for diffusion of soluteatoms in intermetallic compound) willcome into the picture. A further complex situation will arise when multiple layers of different intermetallic compounds areformed. The latter situations were not con-sidered in this work.

When aluminum is to be diffusion welded, E Eff is further increased due to thehindrance caused by aluminum oxide filmpresent on its surface. The film is tena-cious, stoichiometric, and highly stable innature.

In general, during diffusion welding, in-timate contact at the interface is impor-tant; a small contact area leads to a poor

joint. Pressure during diffusion weldinghelps in deformation of contacting asperi-ties at the interface by causing yielding and

Table 1 — Composition and Crystal Structures of Intermetallic Compounds of Fe-Al

Type of Intermetallic Compound Wt-% of Fe Pearson Symbol Crystal Structure

Fe3 Al 86.06 cF16 FCCFeAl 67.31 cP8 CubicFeAl2 50.72 aP18 Triclinic-anorthicFe2 Al5 45.16 mP22 MonoclinicFeAl3 40.70 mC102 BC monoclinic

FeAl6 metastable 25.55 oC28 OrthorhombicFe2 Al9 metastable 31.39 mP22, D8d MonoclinicFeAl x metastable Unknown Unknown Unknown

Table 2 — Mechanical Properties of Cast Fe-Al Intermetallic Compounds

Type of Intermetallic Vickers Hardness Fracture Toughness Compressive CompressiveCompound (9.8 N) K IC, MPa·m1/2 Strength, MPa Strain, %

FeAl3 892 2.15 200 0.00Fe2 Al5 1013 2.30 240 0.00FeAl2 Unknown Unknown Unknown UnknownFeAl 470 Unknown 670 0.45

Fe3 Al 330 Unknown 560 0.80

Fig. 1 — Fe-Al equilibrium diagram. Fig. 2 — Effect of temperature on diffusion coefficients for iron and aluminum.

Fig. 3 — A —Schematic diagram of laser roll welding process; B — details of sheet clamping with lap joint configuration.

A B



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creep deformation (Ref. 11). This makesan intimate contact over a large fraction of the interface area. When high pressuresare applied, not only deformation of as-perities takes place but the aluminum

oxide film is also broken, improving the in-terfacial contact area (Ref. 12). Thus,pressure helps in reducing the effective ac-tivation energy E Eff , which is increased dueto the presence of aluminum oxide film.The values of ESol and kSol will remain con-stant even under varying pressures, be-cause simple diffusion (without formationof intermetallic compound) is involved.

Effect of Interface Layer onJoint Properties

Interface layer thickness has a signifi-cant effect on joint strength: the thinner

the layer, the better the strength. Joints were made between steel and aluminumby deformation welding followed by heattreatment (Ref. 13). High values of frac-ture toughness, GIC, were noted (6097N/m), when the diameter of the Al-richFe2 Al5 was less than 4 µm. The main rea-son for embrittlement of these joints wasattributed to the formation of Kirkendallporosity rather than the intermetalliccompound layer. In the case of resistance

spot welding of mild steel and aluminum Alloy 5052 with aluminum clad steel as aninsert, an increase in shear strength wasobserved from 26 to 56 MPa, when the in-terface layer thickness was reduced from1.0 to 0.4 µm (Ref. 6).

Vacuum roll welded joints of mild steelto aluminum Alloy 5083 showed constantshear strength of 60 MPa when the totalreduction of sheets was above 5% (Ref.14). However, shear strength decreasedafter postheat treatment, as intermetalliccompound layer appeared at the interface.

Solid-Liquid Interaction and Use ofHigh-Energy Beams

When joining temperatures are higherthan the melting point of aluminum, theprocess involves solid steel-liquid alu-minum interaction. During this type of in-teraction, three stages are involved (Ref. 9):∑Wetting of solid steel by molten alu-

minum,∑ Dissolution of iron into liquid alu-

minum,∑Subsequent diffusion of iron in liquid

aluminum.These stages have a strong effect on the

structure and properties of the welded joints. In solid-liquid interaction, formationof intermetallic compounds is considerablyfaster than in solid-solid interaction be-cause atoms have more kinetic energy in theliquid state than in the solid state. However,spreading of liquid aluminum is necessaryfor effective dissolution and diffusion of iron atoms in it. It was suggested that satis-factory spreading is possible when a sharp,intensified, fast moving heat source such asan electron beam is used (Ref. 9).

The high energy density of such asource will melt aluminum quickly, but the

molten pool will be restricted to a smallsize. There will be accelerated diffusion of iron atoms in the molten pool of alu-minum. Fast motion of the heat source willcontrol contact time between solid steeland liquid aluminum and provide a fastcooling rate. This will restrict diffusion toa limited depth and a thinner intermetal-lic compound layer will result, which is fa-

vorable for improving joint strength. If dissolution and diffusion of iron into liq-

uid aluminum are intensified in this way,iron content may be increased beyond 45at-% Fe. Such condition will promote for-mation of Fe-rich intermetallic com-pounds at the interface and will result inimproved joint strength.

A laser beam provides an excellent toolfor achieving control over heat input andin reducing the melt pool size (Ref. 15).

An Nd:YAG laser has been used with and without Al-12%Si welding wire for the joining of aluminum and steel sheets inoverlap configuration (Refs. 15, 16). Theiraverage tensile strength was 171 MPa,

with 3 µm interface layer thickness. Thehigh tensile strength can be attributed tohigh silicon content because the presenceof silicon restricts formation of brittle Al-rich intermetallic compound (Ref. 17).

Fluxes are generally used in brazingand soldering to remove oxides from thebase metal surface by a reduction and/ordissolution reaction, to protect the cleansurface from reoxidation, and to modifythe surface tension of the molten metal(Ref. 18). These functions are helpful dur-ing joining of steel and aluminum. Hence,aluminum brazing flux was used in thepresent work.

The laser roll welding process, whichconsists of simultaneous laser heating andpressing with a roll, has been suggested bythe authors in previous work (Refs. 19–22)to join steel and aluminum. The effect of roll pressure in laser roll welding was inves-tigated previously (Ref. 19). It was notedthat, even with low heat input, increased rollpressure improves the contact area andhelps in intermetallic compound layer for-mation. The effect is similar to that found indiffusion welding in which pressure is an im-portant parameter. When diffusion tem-perature was constant and pressure was in-

Fig. 4 — Effect of holding temperature on the interface microstructures of diffusion welded specimens. A — Holding temperature 823 K; B — holding temperature 873 K.

Fig. 5 — EPMA traces of Fe and Al across the diffusion welded joint showing uniform layer and mixed layer.

A B



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creased, incubation time for formation of the interface layer was reduced. With con-

stant diffusion time, an increase in pressurecaused formation of the interface layer atlower temperatures. Hence, in bothprocesses, pressure plays an important rolein interface layer formation.

With this background, experiments were carried out to join low-carbon steeland aluminum Alloy 5052 by diffusion

welding and laser roll welding. Diffusion welding was carried out by varying tem-perature and time without any insert ma-terial. It was decided to understand the ef-fect of pressure on the kinetics of intermetallic compound layer formationduring diffusion welding. Activation en-

ergy for formation of intermetallic com-pound ( E IMC) and effective activation en-ergy ( E Eff ) values were estimated fordifferent pressures. Time-temperature-phase (TTP) diagrams were plotted fordifferent pressures.

The object of the laser roll welding ex-periment was to promote formation of slightly ductile Fe-rich intermetallic com-pounds in the interface layer and therebyimprove the joint strength by varying laserheat input and roll pressure. Aluminumbrazing flux was applied on the faying sur-face of the aluminum sheet.

Experimental Procedure

Materials Used

Cold rolled, low-carbon steel sheet with 0.12 wt-% carbon and aluminum Alloy 5052-O sheet were used for joining.For diffusion welding, the thickness of both sheets was 1 mm, whereas for thelaser roll welding experiment, 0.5-mm-thick steel sheet and 1-mm-thick alu-minum alloy sheet were used.

Diffusion Welding

From Fig. 2, it can be noted that above723 K (450°C), the diffusion of iron in alu-minum is considerably fast. Liquidus andsolidus temperatures of 5052 aluminumalloy are 922 K (649°C) and 880 K (607°C),respectively. Therefore, diffusion welding

was performed with each permutation of temperature, time, and pressure in thematrix of 773, 798, 823, 848, and 873 K; 2,10, 100, 500, 1000, and 3600 s; and 15.75,31.5, and 47.25 MPa.

Generally, diffusion welding is per-formed by placing sheets of specimen be-tween two punches in a die and the wholeassembly is heated externally. This is doneto keep pressure constant throughout the

thermal cycle. When diffusion welding of steel-aluminum is made in this way, for-

mation of a brittle intermetallic com-pound layer at the interface makes the weld too fragile for handling and furthercharacterization becomes difficult. There-fore, to ensure that specimens do notbreak during characterization, diffusion

welding was conducted in such a way thatthe sheets are not detached from eachother in all stages of welding as well ascharacterization. This was achieved byclamping the sheets with washers, a nut,and a bolt (Material AISI-304). Pressure

was applied with a torque meter prior todiffusion welding and was calculated byusing a formula for threading-friction be-

tween nut and bolt (Ref. 23).Generally, this method suffers from

pressure variation during heating due tothermal expansion. However, pressure re-mains almost constant during diffusion

welding of steel and aluminum, providedthe ratio of their thickness is matched withthat of their linear coefficients of thermalexpansion (a). The values of a for low-carbon steel, 5052 aluminum alloy, andSAE 304 stainless steel are 12.0¥ 10–6, 23.8¥ 10–6, and 18.0 ¥ 10–6K –1, respectively.Thickness of the steel and aluminumsheets are 1 mm each, and they cover a bolt

Fig. 6 — Interface microstructures of diffusion welded specimens with different holding times. A — Holding time 100 s; B — holding time 1 h.

Fig. 8 — Effect of pressure on interface layer mi-crostructure of diffusion welded specimen. A —

Pressure 15.75 MPa; B — pressure 47.25 MPa.

Fig. 7 — Laser beam depth analysis of diffusion welded specimen showing microvoids in interface on steel side. Void locations are shown by arrows.

A A

B

B



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length of 2 mm. Therefore, the change inthe length of the bolt is compensated for bythe change in thickness of the sheets, i.e.,DL Bolt≈ (DL Steel +DL Al) due to heating. For

applied pressures of 15.75, 31.5, and 47.25MPa at room temperature, estimation re- veals they will reduce by 0.02–0.07% in thediffusion temperature range of 773–873 K (500–600°C). This variation is negligibleand therefore can be ignored.

Sheets of steel and aluminum alloy were cut to square coupons of 30 mm perside with an 8-mm-diameter center hole.Faying surfaces were polished with 200grit size emery paper followed by ultra-sonic cleaning in acetone for 5 min. Thesteel and aluminum coupons wereclamped and pressure was applied with atorque meter. A muffle furnace was

heated to a predetermined diffusion tem-perature and soaked for 15 min for tem-perature equalization inside the chamber.The thermal cycle for diffusion welding

consisted of three stages: heating, diffu-sion, and cooling without any protectiveatmosphere. The heating stage was startedby charging a specimen in the furnace.Heating rate was controlled so that diffu-sion temperature was reached within 20 to30 s after charging. This was followed bythe diffusion stage, which consisted of holding the specimen for a predeterminedtime at the diffusion temperature. Afterdiffusion, the specimen was quenched in

water.Diffusion welded specimens were

mounted in self-setting epoxy resin in theas-clamped condition. The mounted spec-

imens were cut along the bolt axis at thecenter, ensuring support to the weldedsheets while revealing a cross section of the joint on either side of the bolt. Thespecimens were polished and etched in3% nital for microstructure observation.Maximum thickness of the interface layerformed was measured for each specimen.Corresponding to the compressive pres-sures applied, separate TTP diagrams

were plotted with diffusion temperatureas ordinate and diffusion time as abscissa.The points where intermetallic compoundlayer thickness was just above 0 µm were

joined with a line to show the “start” curverepresenting incubation times for inter-metallic compound formation.

The laser microscopy technique wasused for depth analysis to detect pore for-mation in and around the interface layer.To identify intermetallic compounds pre-sent in the layer, EPMA by wavelengthdispersive X-ray spectroscopy (WDS) wascarried out on the nital-etched specimen.

Microhardness was measured across theinterface at a 50-g load.

Laser Roll Welding

A schematic diagram of the laser roll welding process is shown in Fig. 3A. A 2.4-kW, continuous-wave CO2 laser facility

was assembled with a flat bending mirrorand roll fixture. The roll is made from AISI304 stainless steel and mounted with a cal-ibrated compression spring for applyingpredetermined roll pressure. Roll pressure

was varied between 150 and 202 MPa. A

quasi-Gaussian laser beam (TEM 01*mode) passes through a focusing lens of ZnSe. A flat mirror placed at 57.5 deg withthe horizontal reflects the laser beam sothat the distance between roll axis and cen-ter of the beam spot is reduced to 17 mm.Shape of the beam was a quasi-ellipticalspot of 2.5 mm minor diameter across and3.5 mm major diameter along the directionof table travel. A defocusing distance of –25 mm was used. Details of the clampingsheets on the table are shown in Fig. 3B.The steel sheet is clamped on top of thealuminum alloy sheet with an overlapping

width of 3 mm and a gap of 0.2 mm.Suitable range of laser power and

travel speed for laser roll welding was de-termined by measuring the maximumtemperature at the interface. The temper-ature was also calculated by using the

Ashby-Easterling model with conductivemode of heat transfer (Refs. 24, 25).

The thickness of the steel sheet was 0.5mm and that of the 5052 aluminum alloy

was 1.0 mm. Sheets were cut in rectangu-lar pieces with 150 mm length and 45 mm

width. After polishing of faying surfaces with a wire brush, both sheets werecleaned in an ultrasonic bath of acetone

Fig. 9 — Microstructure of diffusion welded low-carbon steel-A5052 alloy joint.

Fig. 10 — Stages in formation of interface layer during diffusion welding of low-carbon steel-A5052 alloy sheets

Fig. 11 — Time-temperature-phase (TTP) diagram for diffusion welding of low-carbon steel- 5052 Al alloy with a pressure of 31.5 MPa.

Fig. 12 — Effect of pressure during diffusion welding on the start curves for interface layer formation.



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for 5 min. The steel surface exposed to thelaser beam was coated with graphite to in-crease absorption of the laser beam. To fa-cilitate fast melting of the aluminum alloysheet and to remove the oxide film on it,the faying surface was coated with alu-minum brazing flux, KAlF4: K 2 AlF5.H2O(17–25 wt-%) with particle size of 15 to 21

µm. Argon gas with a 25.0 L/min flow rate was used for shielding. Laser power of 1.5

kW was used for all combinations of travelspeeds and roll pressures.

Laser roll welded specimens were cutacross the lap joint seam for macrostruc-ture and microstructure observation.Etching with 3% nital was made to revealthe interface layer. The thickness of the in-terface layer was measured at five equallyspaced locations and its average value re-ported. Electron-probe microanalysis byWDS was used to analyze the interfacelayer and to identify the intermetalliccompounds present. These results showedlinear traces of iron and aluminum con-

tents in wt-%. Based on the compositionand mechanical properties shown in Ta-bles 1 and 2, thickness of the interfacelayer was divided into Fe-rich intermetal-lic compound thickness (FeAl + Fe3 Al)and Al-rich intermetallic compound thick-ness (FeAl3 + Fe2 Al5). For the tensileshear test, three samples were cut as perthe JIS Z 3192–3B specification with 8 mm

width across the lap joint. The tensi leshear test was carried out on a universaltesting machine.

The sheared surfaces of steel sheet were subjected to X-ray diffraction analy-

sis by using Mo-K aradiation to identify in-termetallic compounds present on it.Backscattered electron microscopy wasused to observe the remaining aluminumarea on the sheared steel surface.

Experimental Results andDiscussions

Diffusion Welding

Diffusion welding of similar materialsgoes through the stages of asperity con-tact, their deformation, formation of in-terfacial boundary, grain boundary migra-tion, and volume diffusion with poreelimination (Ref. 11). However, when dis-similar metals are to be joined, the ab-sence or presence of intermetallic com-pounds on their equilibrium diagramshould be carefully observed. When suchintermetallic compounds exist, then theirdiffusion welding will go through two ad-ditional stages: formation of supersatu-rated solid solution and its transformationinto intermetallic compound (Ref. 10).Such course of intermetallic compoundformation takes place when steel and alu-minum are diffusion welded.

Microstructure of Interface Layer

and Its Development

All intermetallic compound layers ap-peared white in color after etching withnital. Therefore, the type of intermetalliccompounds present cannot be distin-guished microscopically.

Figure 4 shows microstructures and re-spective sketches of diffusion-weldedspecimens at different temperatures witha diffusion time of 10 s and pressure of 47.25 MPa. Formation of the intermetalliccompound layer is shown in Fig. 4A at adiffusion temperature of 823 K (550°C).

The intermetallic compound has juststarted to grow as separate particles (13 µm thickness) at the aluminum side of theinterface (i.e., above the initial interfaceline in Fig. 5A). Such formation in the alu-minum side is in agreement with the fasterdiffusion of iron in aluminum rather than

vice versa.In diffusion welding at 873 K (600°C)

(Fig. 4B), the intermetallic compoundlayer was continuous along the interfaceand it was thicker (41 µm) than thatformed at 823 K (550°C). This layer con-sists of two distinct parts, a “uniform”layer (U) and a “mixed” layer (M). Theuniform layer has a spike-shaped bound-ary protruding in the steel side and it con-tains some black lines perpendicular tothe interface. As seen in the EPMA tracesof Fig. 5, the uniform layer mostly consistsof FeAl3 with 40.7 wt-% Fe. The mixedlayer is seen on the aluminum side, and itconsists of white islands surrounded by agray phase. The EPMA trace of aluminumin the mixed layer shows wide variations.The peaks correspond to the white islandsof aluminum and the surrounding grayphase is made of the Al-rich intermetalliccompound FeAl3.

The effect of holding time on the mi-crostructure of the interface layer is shownin Fig. 6. When holding time was 100 s,temperature 823 K (550°C), and pressure31.5 MPa (Fig. 6A), the intermetallic com-pound layer thickness was about 9 µm.With a holding time of 3600 s and theother conditions being the same, the layerthickness increased to 46 µm — Fig. 6B.When holding time was 100 s, particles of intermetallic compound are seen formedand spread partially along the interface.However, they are not yet seen merged

Fig. 13 — Effect of pressure on activation energy for formation of intermetallic compound layer at the interface of diffusion welded low-carbon steel-5052 aluminium alloy joints.

Fig. 14 — Maximum temperatures at the interface when laser heated with different travel speeds and laser powers. Calculated temperatures are shown by lines and measured temperatures are shown as leg- ends.

Fig. 15 — Effect of travel speed on macrostruc-tures of laser roll welded specimen .A — Without

laser heating; B — 1.5 m/min; C — 1.8 m/min; D — 2.2 m/min.

A

B

C

D



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with the neighboring ones — Fig. 6A. Inthe case of the specimen with a 1-h hold-ing time (Fig. 6B), the interface mi-crostructure shows a uniform layer withspiked boundary on the steel side and themixed layer is almost absent. The EPMA results show that long diffusion timecauses transformation of aluminum is-lands into FeAl3 intermetallic compoundand, hence, the uniform layer grows at theexpense of the mixed layer.

Fine spike lines are seen perpendicularto the interface line in the uniform layeron the steel side (Figs. 4B and 6B). Thereis considerable migration of iron towardthe aluminum side along the grain bound-aries and through grains, whereas migra-tion of aluminum is negligible. Therefore,deficiency of iron atoms occurs at someplaces in the steel side of the interface.

Particularly, migration of iron atoms fromgrain boundaries and lattice defects is fastas compared to that from the grain bulk.Therefore, deficiency of iron atoms andpoor supply of aluminum atoms results inaccumulation of microvoids at the grainboundaries and lattice defects. These mi-crovoids grow in size and ultimately formlinear voids in the uniform layer. Theselines are perpendicular to the interfaceand indicate strong flow of iron atoms to-

ward the aluminum side. Hence, it isthought that the Kirkendall effect is thereason for linear void formation in theuniform layer.

Depth analysis (on laser microscope)of the diffusion-welded interface has beenreported in previous work (Ref. 19). Itconfirms the existence of voids at the spikelines in the uniform layer — Fig. 7. Mag-

nified microstructure at the steel side of the interface layer is shown on top andcorresponding depth values are shown inthe bottom graph. Depth values at loca-tions marked with arrows are in the rangeof 0.7–2.1 µm. These locations correspondto the black lines in the uniform layer.They are deeper than the etched grainboundaries in steel. In the same photo-graph (Fig. 7), the boundary between steel

and intermetallic compound layer is seenas a spiked structure with rounded tips. Itis thought that these rounded tips areformed due to diffusion of iron not onlyfrom the grain boundaries but also fromthe bulk grains.

Figure 8 shows the effect of appliedpressure on the microstructure of diffu-sion-welded joints. The conditions werediffusion temperature 873 K (600°C),holding time 100 s, and pressures 15.75and 47.25 MPa. When pressure was 15.75MPa, layer thickness measured 37 µm(Fig. 8A), whereas with 47.25 MPa, it mea-

sured 59 µm — Fig. 8B. The effect of pres-sure is also seen on the shape of ferritegrains in steel. Equiaxed grains are seen

when pressure is low (15.75 MPa) andelongated grains are seen when pressureincreased (47.25 MPa). The uniform layeris thicker and the spiked boundary iscoarser in the case of high pressure thanthat in the case of low pressure. Mixed lay-ers are clearly seen in both microstruc-tures. With high pressure, it is thought thatthe amount of contact area for diffusionincreases and helps in accelerating migra-tion of atoms across the interface.

Figure 9 shows the complete mi-crostructure of the intermetallic com-pound layer formed along the interface. It

was diffusion welded at 873 K (600°C) for500 s with 15.75 MPa and shows a maxi-mum thickness of about 120 µm. If thethickness is divided into two parts at theoriginal interface line, the thicker part isseen in the aluminum side (about 100 µm)and the thinner one in the steel side (about20 µm).

From observing the changes in the mi-crostructures with diffusion time, stages inthe intermetallic compound layer forma-tion can be understood. They are schemat-ically shown in Fig. 10, as follows:∑ Formation of separate particles of

intermetallic compounds in the aluminumside.∑ Particles are connected along the in-

terface; start of uniform layer formation.∑ Mixed layer formation starts in the

aluminum side.∑ The mixed layer grows in the alu-

minum side and the uniform layer growsslightly in the steel side with spiked bound-ary.∑ Growth of the uniform layer at the

expense of the mixed layer.

Fig. 16 — Microstructures of laser roll welded specimen showing effect of travel speed. A — 1.8 m/min; B — 2.0 m/min; C — 2.4 m/min.

Fig. 17 — Electron-probe microanalysis traces of Fe and Al across laser roll welded interface layer. A—Travel speed 1.2 m/min; B — travel speed 2.0 m/min.

A

A

B

B

C



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Time-Temperature-Phase Diagram

The TTP diagram for pressure of 31.5MPa is shown in Fig. 11. The curved lineson the diagram show thickness of the in-terface layer formed with different combi-nations of diffusion temperature and time.The line showing 1 µm thickness is the“start” curve for formation of the inter-face layer. The region below the start

curve indicates the “incubation time” dur-ing which diffusion continues to form su-persaturated solid solution.

The effect of a change in pressure onstart curves is shown in Fig. 12. During theearly stages of diffusion welding at hightemperatures, increasing pressure causesstart curves to shift downward, indicatingshorter incubation times and lower temper-atures for formation of the interface layer

with the same thickness. No incubation time was observed for high temperatures of dif-fusion, which means diffusion starts duringthe heating stage itself, indicating rapid

atomic movement at high temperatures.High pressures and high temperatures

accelerate intermetallic compound for-mation. There is the combined effect of in-creased contact area due to pressure andincreased atomic mobility due to hightemperature. When pressure is increased,contact area increases due to hastening of the initial stages of diffusion, namely sur-face deformation and asperity collapse(Ref. 11). In addition to increased surfacedeformation, an increase in pressure alsohelps in breaking of the aluminum oxidefilm (Ref. 12).

However, incubation times are longerin the case of low diffusion temperatures.The effect of high pressure on loweringstart curves gradually fades away at lowdiffusion temperatures. At these temper-atures, mobility of atoms is reduced due toinsufficient kinetic energy. Though asper-ity collapse, surface deformation, andbreaking of oxide film are taking place dueto high pressure, longer incubation timesare needed at low diffusion temperatures.

Kinetics for Growth of Intermetallic

Compound Layer

The effect of temperature on kineticsof diffusion is well known and can be ex-pressed by the Arrhenius equation (Equa-tion 2). The value of activation energy ESol

for diffusion of iron atoms in aluminumlattice is 135 kJ/mol (Ref. 7). The effect of pressure applied during diffusion weldingon the effective activation energy E Eff forthe formation of the intermetallic com-pound is quantitatively investigated. The

values of E IMC for different pressures werecalculated from growth rate and Arrhe-nius plots. Slope of the Arrhenius plots is

the activation energy for intermetallic

compound formation.The maximum value of activa-

tion energy for formation of inter-metallic compounds E IMC during dif-fusion welding is 228.9 kJ/mol withan applied pressure of 15.75 MPa.The value of effective activation en-ergy E Eff is given by using Equation 1(93.9 kJ/mol). As explained in theintroduction, E Eff takes into accountthe factors nW , Ecr , and ne (Ref. 10).Factor nW is due to the difference inlocal composition fluctuations inthe aluminum alloy matrix and theintermetallic compound, mostlyFeAl3. A crystal structure changefrom aluminum to monoclinic FeAl3

(Table 1) is responsible for increas-ing the value of Ecr . It also causes de-formation of matrix increasing ne.

Figure 13 shows the effect of pres-sure on activation energy for forma-tion of intermetallic compound E IMC. Acti-

vation energy decreases from 228.9 to 193.6kJ/mol as pressure is increased from 15.75to 47.25 MPa. As explained in the previoussection, high pressure increases the metal-to-metal contact area by surface deforma-tion, asperity collapse, and formation of anew surface by breaking the oxide film onthe aluminum surface. This decreases ef-

fective activation energy E Eff from 93.9 to58.6 kJ/mol, when pressure is increasedfrom 15.75 to 47.25 MPa. Decrease in ef-fective activation energy E Eff means the en-ergy barrier for diffusion and formation of the intermetallic compound is reduced. Inother words, diffusion and intermetalliccompound formation are facilitated by in-creasing pressure.

Temperature Field Simulationfor Steel Sheet

Temperature reached at the interface is

a critical parameter in laser roll welding of

steel to aluminum alloy. For promotingformation of Fe-rich intermetallic com-pounds in the interface layer, two condi-tions are important: a large number of iron atoms — beyond 45 at-% Fe —should diffuse in aluminum in a short pe-riod of time, and formation of brittle in-termetallic compounds should be sup-pressed.

For accelerated diffusion of ironatoms, liquid aluminum is a better phasethan solid and it should wet the steel sur-face sufficiently (Ref. 9). Therefore, theinterface temperature should exceed theliquidus temperature of the aluminumalloy. Because brittle intermetallic com-pounds are formed between 1433 and1442 K (1160 and 1169°C), interface tem-perature should rapidly increase above1473 K (1200°C). It should also decreaserapidly through that temperature range.Laser heating involves not only a rapidheating rate but also rapid cooling, reduc-

ing the time for formation and growth of

Fig. 18 — Effect of travel speed on average interface layer thickness and percentage of Fe-rich and Al-rich intermetallic compounds.

Fig. 19 — Effect of average interface layer thickness on average shear strength of laser roll welded low-carbon steel-5052 aluminium alloy joints.

Fig. 20 — X-ray diffraction pattern on steel surface after tensile shear testing. Existence of ductile FeAl intermetal-

lic compound is shown by arrows.



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JANUARY 2004-S24

brittle intermetallic compounds.To determine laser power and travel

speed suitable to reach this level of tem-perature at the interface, the Ashby-East-erling model involving laser heating in theconduction mode was used (Refs. 24, 25).In this model, the heat source is consid-ered as a moving finite line situated abovethe surface and parallel to it. In the pre-sent work, the laser beam spot is quasi-

elliptical in shape because the incidenceangle is 57.5 deg, reducing the energy ab-sorption. Therefore, an absorptivity of 0.3

was used. Other parameters used for cal-culation of temperature were thermal dif-fusivity 0.19 ¥ 10–4 m2 /s, thermal conduc-tivity 15.1 J/m/s/K, and specific heat per

volume 9.6 ¥ 106 J/m3 /K. Calculated andexperimentally measured temperatures atthe bottom of the steel sheet fairly coin-cide as shown in Fig. 14. From these re-sults, laser power of 1.5 to 2.0 kW wasidentified as suitable for laser roll welding.Hence, for laser roll welding, 1.5-kW-laser

power was used and travel speed was var-ied between 1.0 and 3.0 m/min.

Laser Roll Welding

Macrostructures of Laser Roll Welded Joints

The effect of travel speed onmacrostructures of laser roll welded spec-imens with a roll pressure of 202 MPa andlaser power of 1.5 kW is shown in Fig. 15.To understand the difference inmacrostructures, a photograph of sheets

without laser heating is shown in Fig. 15A.

When a slow travel speed of 1.5 m/min wasused, a considerable change in the originalshape is seen — Fig. 15B. The figureshows melting occurs in both the steel andaluminum sheets because of the highamount of laser heat input. Aluminumspreading takes place along the bottom of the steel sheet. Though welding had takenplace, it was not uniform. When travelspeed is increased to 1.8 m/min, melting of steel in a shallow depth is seen on the topsurface. However, the aluminum sheet hasundergone considerable melting andspreading — Fig. 15C. With a fast travelspeed of 2.2 m/min, the macrostructureshows no melting of steel and aluminummelting and its spreading are limited —Fig. 15D. A similar effect of travel speedon macrostructure was observed with 150and 175 MPa roll pressures.

Laser heating provides high energydensity; therefore, temperature of thesteel sheet at the interface rapidly in-creases above 1473 K (1200°C). When thehot surface of the steel comes in contact

with the aluminum due to roll pressure, itmelts and spreads along the bottom sur-face of the steel sheet. Presence of alu-minum brazing flux must have supported

in these stages. This situation leads tosolid-liquid interaction and acceleratesiron diffusion in molten aluminum.

Microstructure and Thickness

of Interface Layer

Interface microstructures of laser roll welded joints with 175-MPa roll pressureand 1.5-kW laser power are shown in Fig.

16. In contrast to the microstructures pro-duced by diffusion welding, microstruc-tures of the laser roll welded specimenshowed only the “uniform” layer, irre-spective of travel speed and roll pressure.The “mixed” layer is completely missing.For a slow travel speed of 1.8 m/min and175-MPa roll pressure, the average layerthickness is 8 µm — Fig. 16A. A spikedboundary at the steel side of the interfacelayer is seen similar to that observed in dif-fusion-welded joints. With the slow travelspeed, the Kirkendall effect is active informing microvoids along the spiked

boundary; but they are fine due signifi-cantly to the short times for diffusion.

For the intermediate travel speeds, theaverage interface layer thickness is re-duced sharply. When travel speed is 2.0m/min, the interface thickness is 5 µm andthe spiked boundary is not seen — Fig.16B. When travel speed is increased fur-ther to 2.4 m/min, the interface layer fur-ther reduces to 4 µm without spikedboundary — Fig. 16C. For all roll pres-sures used, a similar effect of reduction ininterface layer thickness with increasingtravel speed was observed.

The backscattered electron imageshowed a dendritic structure in the heat-affected zone of aluminum adjacent to theinterface layer, confirming aluminummelting during laser roll welding.

Effect of Travel Speed on Composition of

Interface Layer

As the optical microstructure does notreveal the details of the interface layer,EPMA of iron and aluminum across it

were made to identify intermetallic com-pounds. Two such traces for specimens

welded with a pressure of 150 MPa andtravel speeds of 1.2 and 2.0 m/min areshown in Fig. 17. Intermetallic com-pounds in the interface layer are shown indark bold letters and black arrows. Faintgray letters and arrows are used for thoseintermetallic compounds not present inthe interface layer.

Variation in composition of the inter-metallic compound layer is seen asstepped lines. Some intermetallic com-pounds are very narrow in thickness andare closely spaced together. Therefore,overlapping of electron beam on neigh-boring intermetallic compounds gives in-

termediate composition, which is inaccu-rate. Sometimes presence of complex mixed structure may also yield an inaccu-rate result. When intermetallic com-pounds are very thin, as seen in Fig. 17B,it is difficult to estimate exact composi-tion.

With a slow travel speed of 1.2 m/min,FeAl3 is predominantly observed on thealuminum side with some Fe2 Al5 at the

center — Fig. 17A. On the steel side, thelayer mainly consists of FeAl with someFe3 Al. For a faster travel speed of 2.0m/min, the interface layer has becomethinner and the thickness of FeAl3 is fur-ther reduced — Fig. 17B. From Table 2, itcan be noted that FeAl3 and Fe2 Al5 arehard and brittle, whereas FeAl and Fe3 Alare slightly ductile. Hence, the interfacelayer is divided into a brittle part (FeAl3 +Fe2 Al5) and a slightly ductile part (FeAl +Fe3 Al), based on the composition given bythe EPMA traces. The fraction of theseparts is expressed in percentage with re-

spect to total thickness of interface layer.Figure 18 shows the effect of travel

speed on average interface layer thicknessand the percentages of the Al-rich(FeAl3+Fe2 Al5) and Fe-rich parts(FeAl+Fe3 Al) in it, as measured from theEPMA results. Changes taking place witha roll pressure of 175 MPa are shown inFig. 18. Average interface layer thicknessdecreases sharply from 14 to 8 µm whenthe travel speed increases from 1.5 to 1.8m/min; but the percentage thickness of the

Al-rich part reduces slightly from 76 to75%. As the travel speed increases in the

intermediate range of 1.8–2.0 m/min, av-erage interface layer thickness decreasesslowly from 8 to 5 µm. In this intermediaterange, the thickness of the Al-rich inter-metallic compounds decreases from 75 to70%. For a higher travel speed of 2.4m/min, total interface layer thickness isreduced further to 4 µm and the percent-age of Al-rich intermetallic compound de-creases to 50%. As the travel speed in-creases, percentage of Fe-richintermetallic compound increases at theexpense of the Al-rich intermetallic com-pounds. Similar effect of travel speed oncomposition of interface layer was seenfor roll pressures of 150 and 202 MPa.

Laser heating and the presence of alu-minum brazing flux have caused rapidmelting and spreading of aluminum above1473 K (1200°C). Simultaneously, diffu-sion of a large number of iron atoms in liq-uid aluminum might have caused forma-tion of Fe-rich intermetallic compound onthe steel side of the interface (Ref. 9). Thethermal cycle of laser heating provides acondition that is far away from equilib-rium, and time for formation of brittle Al-rich intermetallic compounds on the steelside is not sufficient. Therefore, formation



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of brittle Al-rich intermetallic compoundsis suppressed and that of the Fe-rich in-termetallic compounds is promoted, atleast on the steel side of the interfacelayer. Electron-probe microanalysis re-sults showed that complete elimination of

Al-rich brittle intermetallic compounds isnot taking place. They are always formedon the aluminum side of the interfacelayer. This may be attributed to the pres-

ence of the large amount of aluminum andto their low free energy values (Ref. 4).

Tensile Shear Test Results

When a lap joint of dissimilar metals issubjected to tensile shear testing, the spec-imen may fail either at the interface or inthe weaker base metal. When failure oc-curs at the interface, the shear strength of the joint is calculated based on the overlaparea. If the specimen fails in the weakerbase metal, it always fails in the tensilemode. Failure of the weaker base metal in

tension indicates that the shear strength of the joint is greater than the tensilestrength of the weaker base metal. Whenthe weaker base metal fails, joining of thedissimilar metals can be regarded as suc-cessful. However, the exact value of the

joint’s shear strength remains unknown.The tensile shear test specimen has a

width of 8.0 mm. Tensile strength of the5052-O aluminum alloy is 200 MPa and itscross-sectional area 8.0 mm2 (with a thick-ness of 1.0 mm), hence it will support amaximum tensile force of 1.6 kN. The ten-sile strength of low-carbon steel is 300

MPa and its cross-sectional area is 4.0 mm2

(with a thickness of 0.5 mm); hence, it willsupport a maximum tensile force of 1.2kN. Therefore, the steel side of the joint

will fail first due to its smaller cross section(even though it has a higher tensilestrength), whereas the cross section of thealuminum alloy is twice that of the steeland will therefore have a greater load-carrying capacity (despite its lower yieldstrength). The steel side will fail if the jointhas a higher shear strength than 50 MPa asshear area is 24 mm2 (8 mm wide ¥ 3 mmoverlap length).

With slow travel speeds, there is exces-sive laser heat input and the cooling rate isslow. This provides sufficient time for dif-fusion and formation of a thick interfacelayer containing a large amount of Al-richbrittle intermetallic compounds. Such aninterface structure drastically reduces the

joint strength. When travel speeds arefast, there is insufficient heat and time formelting of the aluminum and the diffusionprocess to take place. This results in in-complete welding at the interface andpoor shear strength. Intermediate travelspeeds provide an adequate amount of heat and cooling rate, so that a thin con-

tinuous interface layer is formed with a re-duced amount of brittle Al-rich inter-metallic compound — Fig. 18.

When the average interface layerthickness is between 4 and 5 µm with thepercentage of Fe-rich intermetallic com-pounds between 25 and 40%, shearstrength reaches a maximum value be-tween 50.8 and 55.9 MPa — Fig. 19. Thesespecimens failed partially — at the inter-

face and in aluminum. The rest of thespecimens — belonging to all combina-tions of roll pressure and travel speed —failed at the interface. This indicates thatthe joints have a lower shear strength thanthe tensile strength of the low-carbon steelsheet, which has a lower load-carrying ca-pacity in the present combination. Thereason for this may be attributed to thelarge amount of Al-rich brittle intermetal-lic compounds (60–75%) still present inthe interface layer. However, promotionof Fe-rich intermetallic compounds in thelayer has improved the shear strength

from 11.0 to 55.9 MPa, which is shown asa gray area in Fig. 19. As the intermetalliccompound thickness decreases, theamount of Fe-rich intermetallic com-pounds in it increases and that of Al-richintermetallic compound decreases —Figs. 18, 19.

X-ray diffraction pattern on steel sur-face after shear testing is shown in Fig. 20.It shows that the Fe-rich intermetalliccompound FeAl exists in the interfacelayer in addition to Al-rich intermetalliccompounds FeAl3 and Fe2 Al5. This con-firms the promotion of Fe-rich inter-

metallic compound formation in the inter-face layer by laser roll welding. A backscattered electron image of thesheared steel surface — belonging to thecondition 150 MPa roll pressure and 1.6m/min travel speed — showed that about70% area is covered with aluminum.

Conclusions

The present study is focused on joininga dissimilar metal combination of low-carbon steel and 5052 aluminum alloy.Diffusion welding was studied for the ef-fect of pressure on kinetics of interfacelayer formation. Laser roll welding was in-

vestigated to achieve a thin interface layerthickness and to promote a slightly ductileFe-rich intermetallic compound in it to in-crease the joint strength. The followingconclusions can be drawn.

1) The diffusion welded interface layercontains only the brittle intermetalliccompounds FeAl3 and Fe2 Al5. It can be di-

vided into two parts: a uniform layer onthe steel side and a mixed layer on the alu-minum side. The uniform layer consists of FeAl3 compound and protrudes in thesteel side as a spiked boundary. It also con-

tains fine linear voids, which are thoughtto be the result of the Kirkendall effect.The mixed layer on the aluminum sideconsists of a mixture of FeAl3 compoundand aluminum.

2) Activation energy for formation of intermetallic compound ( E IMC) is greaterthan that of diffusion of atoms in metals( ESol). During diffusion welding of low-carbon steel/5052 aluminum alloy, as pres-

sure is increased from 15.75 to 47.25 MPa,the value of ESol does not change (135kJ/mol), whereas that of E IMC decreasesfrom 228.9 to 193.6 kJ/mol (with corre-sponding decrease in E Eff ). This indicatesthat intermetallic compound formation isfacilitated by increasing pressure.

3) Laser roll welded joints showedmelting and spreading of the aluminumalloy on the bottom surface of the steelsheet. This may be attributed to the highenergy density of the laser beam and thepresence of flux at the interface. Suchmelting and spreading might have led to

liquid aluminum-solid iron interaction atthe interface, accelerating diffusion of iron in liquid aluminum, and promotingslightly ductile Fe-rich intermetallic com-pound formation. Electron-probe micro-analysis results and X-ray diffractionanalysis confirmed the existence of Fe-richintermetallic compounds on the steel side.

4) The effect of roll pressure in laserroll welding was similar to that observed indiffusion welding. It increased the contactarea and effectively facilitated diffusion ina short time. Increase in travel speed ledto thinning of the interface layer due to re-

duced laser heat input. Thinning down of the interface layer is also accompanied byan increase in the percentage of Fe-rich in-termetallic compounds (FeAl+Fe3 Al) atthe expense of Al-rich intermetallic com-pounds (FeAl3+Fe2 Al5). However, com-plete elimination of brittle intermetalliccompounds is not possible due to the largeamount of aluminum on the aluminumside of the interface.

5) In the diffusion welded interfacelayer, Fe-rich compounds were not seen.Laser roll welding achieved their promo-tion from as low as 23% to as high as 51%.This change in composition of the inter-face layer improved the shear strength of

joints from 11.0 to 55.9 MPa. When the av-erage interface layer thickness is between4 and 5 µm, with the percentage of Fe-richintermetallic compounds between 25 and40%, maximum shear strength is reachedbetween 50.8 and 55.9 MPa.

6) Except for the specimens with shearstrength higher than 50 MPa, all laser roll-

welded specimens failed at the interface.Therefore, the shear strength of the jointsis lower than the tensile strength of thelow-carbon steel sheet, which has a lowerload-carrying capacity in the present com-



WELDING RESEARCH

JANUARY 2004-S26

bination. This may be attributed to thelarge amount of Al-rich brittle intermetal-lic compounds (60–75%) still present inthe interface layer.

Acknowledgment

The authors are thankful to AppliedLaser Engineering Center (ALEC),Japan, for its financial help with this re-

search.

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Documents

Aluminum Alloy 5052 and L-C Steel by Laser Roll Welding