welding journal | january 2013

Introduction

Prized for its excellent strength-to-weight ratio, magnesium and its alloys arecurrently under intense investigation foruse in many applications in the automo-tive and aerospace industries (Refs. 1–3).However, steel sheet is still the most com-monly used material in the automotive in-dustry for fabrication of autobody struc-tures. The ability to make hybridstructures of magnesium alloy and steelsheet would be desirable for many appli-cations in the automotive industry, be-cause the overall weight of the autobodycould be reduced resulting in better fuelefficiencies and lower environmental im-

pact. Therefore, there is increasing inter-est in identifying and developing new tech-niques and processes that can be used tomake dissimilar joints between magne-sium alloys and steel sheet (Refs. 3–9).

Joining magnesium alloys to steel byconventional fusion welding technologiesis difficult due to the large difference inthe melting points between Mg (649°C)and Fe (1538°C). In addition, the boilingpoint of magnesium is only 1091°C, so di-rect contact with molten steel causes cata-strophic vaporization of the magnesium(Refs. 3–8). Moreover, the maximum solid

solubility of Fe in Mg is estimated to beonly 0.00041 at.-% Fe (Ref. 6) and wettingof steel by molten magnesium is very poor(Ref. 8).

The weldability of magnesium to steelusing the hybrid laser-arc welding (Refs. 5,6, 8, 9) and resistance spot welding (Refs.4, 7) processes has been examined. Zhaoet al. (Ref. 5) used a hybrid laser-gas tung-sten arc welding (GTAW) process to joinAZ31B magnesium alloy and 304 stainlesssteel. However, oxides that formed at theinterface were found to cause joints withpoor tensile strength. Using the samewelding technique, Liu et al. (Refs. 8, 9)studied lap joining of AZ31B Mg alloy toQ235 steel with Sn and Cu interlayers.Mg2Sn and Mg2Cu intermetallic com-pounds were found to form along the grainboundaries of the Mg alloy when using theSn and Cu interlayers, respectively. Theuse of Sn and Cu interlayers was reportedas the main reason for the elimination ofgaps along the steel-fusion zone interfaceand the improvement of wetting proper-ties of the steel by molten magnesiumalloy (Refs. 8, 9). Finally, in a recent study,Liu et al. (Refs. 4, 7) used resistance spotwelding to join AZ31B magnesium alloyto DP600 Zn-coated steel. They foundthat a preexisting transition layer ofFe2Al5 between the Zn coating and thesteel improved wetting and bonding be-tween the steel and the magnesium alloy.

Review of the literature suggests thatjoining Mg alloys to steel will be possibleprovided the temperatures required forjoining are kept below the boiling point ofthe magnesium alloy (1091°C) and pro-vided another interlayer element is usedthat can interact and promote wetting andbonding between both immiscible alloys.For this reason, brazing can be a superiorchoice in joining dissimilar metals such asmagnesium and steel because brazing

SUPPLEMENT TO THE WELDING JOURNAL, JANUARY 2013Sponsored by the American Welding Society and the Welding Research Council

Interfacial Microstructure of Diode LaserBrazed AZ31B Magnesium to Steel Sheet

Using a Nickel Interlayer

The formation of a nano-scale Fe(Ni) transition layer on the steel during laserbrazing was found to be responsible for the formation of a metallurgical bond

between the steel and magnesium

BY A. M. NASIRI, D. C. WECKMAN, AND Y. ZHOU

KEYWORDS

Laser BrazingAZ31B Mg SheetSteel SheetDissimilarIntermetallic Compound

A. M. NASIRI([email protected]), D. C.WECKMAN, and Y. ZHOU are with Departmentof Mechanical & Mechatronics Engineering, Cen-tre for Advanced Materials Joining, University ofWaterloo, Waterloo, ON, Canada.

ABSTRACT

The brazeability of AZ31B-H24 magnesium alloy and steel sheet with a microlayerof electro-deposited Ni in a single flare bevel lap joint configuration has been investi-gated. The macro- and microstructure, element distribution, and interfacial phases of thejoints were studied by optical microscopy (OM), scanning electron microscopy (SEM),transmission electron microscopy (TEM), and X-ray diffraction (XRD). Wetting of thesteel by the Mg-Al brazing alloy was improved significantly through the addition of a Nielectroplated interlayer. Bonding between the magnesium brazing alloy and the steel wasfacilitated by the formation of a transition layer composed of a solid solution of Ni in Feon the steel followed by a layer of α-Mg + Mg2Ni eutectic. A band of AlNi intermetal-lic compound with different morphologies also formed along the steel-fusion zone in-terface, but was not directly responsible for bonding. Ni electroplating was found to sig-nificantly improve the brazeability and mechanical performance of the joint. The averagefracture shear strength of the bond reached 96.8 MPa and the joint efficiency was 60%with respect to the AZ31B-H24 Mg alloy base metal. In all cases, failure occurred in thefusion zone very close to the steel-fusion zone interface.

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Nasiri Supplement January 2013_Layout 1 12/12/12 1:47 PM

temperatures are generally lower than themelting points of both base metals. In ad-dition, very fast heating and cooling ratescan be used during the brazing process tominimize the thickness of intermetalliccompounds that might form along the in-terfaces (Ref. 10).

The benefits of using laser brazing andlaser welding-brazing technologies forjoining dissimilar materials are also be-coming increasingly recognized due to thecombined attributes of furnace brazingand laser welding (Ref. 11). With a morelocalized energy input and more precisecontrol of the laser beam energy, highjoining speeds and accompanying highcooling rates can be realized with minimalheating of the parts. Also, laser brazingand laser welding-brazing can prevent orminimize excessive formation of detri-mental intermetallic phases. If intermetal-lic layers can be limited to thicknessesbelow 10 μm, acceptable joint strengthsand mechanical properties may be ob-tained (Refs. 12–14).

In our previous study (Ref. 15), a diodelaser brazing process was developed forjoining Mg alloy sheet to aluminized steelsheet where the Al-12Si coating served asthe interlayer. This coating was found topromote wetting of the steel by the mag-nesium brazing alloy; however, a preexist-ing layer of brittle θ-FeAl3 along thebraze-steel interface was found to degradethe mechanical properties of the joint as

failure of the joint always occurred by frac-ture of this brittle intermetallic layer. Fol-lowing a review of binary and ternaryphase diagrams, nickel was identified as apotentially viable interlayer element be-tween the steel and Mg-9Al-2Zn brazingalloy used. Therefore, the purpose of thispresent study was to investigate the braze-ability, interfacial microstructure, and me-chanical properties of the laser brazedAZ31B-H24 magnesium alloy to steelsheet with an electrodeposited layer of Nion the steel to act as the interlayer ele-ment. It is expected that development ofthis laser brazing technology for joining ofsteel-interlayer-Mg alloy combinationswith a strong metallurgical bond betweenthe steel and Mg alloy will facilitate in-creased application and use of Mg alloysin the automotive industry.

Experimental Procedure

In this study, 2-mm-thick commercial-grade twin-roll strip cast AZ31B-H24 Mgalloy sheet and 1-mm-thick steel sheetwere used as the base materials. Thechemical compositions of the base materi-als are given in Tables 1 and 2. A 2.4-mm-diameter TiBraze Mg 600 filler metal(Mg-Al-Zn alloy) with solidus and liq-uidus temperatures of 445° and 600°C, re-spectively, was chosen for this study. Thecommercial flux used in the experimentswas Superior No. 21 manufactured by Su-

perior Flux and Manufacturing Co. Thispowder flux was composed of LiCl (35–40wt-%), KCl (30–35 wt-%), NaF (10–25 wt-%), NaCl (8–13 wt-%), and ZnCl2 (6–10wt-%) (Ref. 16).

The AZ31B Mg and steel sheets werecut into 60- × 50-mm specimens. Prior tolaser brazing, the oxide layers on the sur-faces of the magnesium sheets were re-moved by stainless steel wire brushing. Allthe specimens were ultrasonically cleanedin acetone to remove oil and other con-taminants from the specimen surfaces.

The edge of each steel sheet was bentin order to make a single-flare bevel lapjoint after clamping against the magne-sium sheet. After bending, the steel speci-mens were cleaned in acetone and thenground to 1000 grit using SiC abrasivepaper and again ultrasonically cleaned inacetone. The prepared surfaces were thenimmediately electroplated with elec-trolytic pure nickel. In the Ni electroplat-ing process, the clean steel sample was thecathode and graphite was the anode. Thecomposition of the electroplating solutionand the electroplating conditions arelisted in Table 3. Figure 1A shows aschematic of the Ni electrodepositionprocess used. In order to get a uniform 5-μm-thick Ni layer on the steel, differentcathode current densities and platingtimes were tested. Electrodeposition of Niusing a cathode current density of 120mA/cm2 for 10 min was found to provide a

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Fig. 1— A — Schematic of the Ni electrodeposition process on steel; B —transverse section of the Ni electrodeposited layer on the steel substrate.

Table 1 — Measured Chemical Composition of the AZ31-H24 Mg Alloy Sheet and TiBraze Mg 600Filler Metal (wt-%)

Al Zn Mn Si Mg

AZ31B-H24 3.02 0.80 0.30 0.01 Bal.TiBraze Mg 600 9.05 1.80 0.18 — Bal.

Table 2 — Measured Chemical Composition ofthe 1-mm-Thick Steel Sheet (wt-%)

C 0.01Mn 0.5P 0.010S 0.005Fe Bal.

A B


5.5 ± 0.9-μm-thick pure Ni coating layeron the steel with a defect-free interface.Figure 1B shows a SEM micrograph of thecross section of the nickel-coated steel.The white layer on top of the steel is theNi coating layer. The coating was of uni-form thickness with a void-free interface.Energy-dispersive X-ray spectrometer(EDS) analysis of the electrodepositedlayer on the steel showed a pure Ni coat-ing layer.

After the electroplating process, theprebent steel sheet was clamped againstthe magnesium sheet to make a single-flare bevel lap joint as shown in Fig. 2A.The filler metal was cut into pieces andpreset on the workpiece at the weld inter-face with some flux before heating andbrazing by the laser beam.

An integrated Panasonic 6-axis robotand Nuvonyx diode laser system with amaximum power of 4.0 kW and a 0.5- ×12-mm rectangular laser beam intensityprofile at the focal point was used for laserbrazing. This energy distribution is moresuitable for brazing processes comparedwith the nonuniform Gaussian-distributedcircular beams generated by CO2 andNd:YAG lasers (Ref. 17). The beam wasfocused on top of the filler metal.

In order to limit oxidation, heliumshielding gas was provided in front of themolten pool with a flow rate of 30 L/minfrom a 6-mm-diameter soft copper feed-ing tube. Laser brazing was performedusing a range of laser powers, travelspeeds, and beam offset positions.

After laser brazing, transverse crosssections of the brazed specimens were cutand mounted in epoxy resin. The sampleswere then mechanically polished using300, 600, 800, 1000, and 1200 grades of SiCgrinding papers followed by polishingusing a 1-μm diamond suspension. Thepolished specimens were etched to revealthe microstructure of the braze metal andAZ31B base material. The etchant was

comprised of 20 mLacetic acid, 3 g picricacid, 50 mL ethanol,and 20 mL water (Ref.18).

Macro- and mi-crostructures of theetched joints were ex-amined using an opticalmetallographic micro-scope. The microstruc-ture and composition ofdifferent zones of thejoint cross section weredetermined using aJEOL JSM-6460 scan-ning electron micro-scope (SEM) and EDS.Phase characterizationof the phases formed inthe steel-fusion zone in-terface and on the frac-ture surfaces was car-ried out using X-raydiffraction (XRD)phase analysis in aRigaku AFC-8 diffrac-tometer with Cu target,50 kV acceleration volt-age, and 40 mA current.

A transmission electron microscope(TEM) foil of the steel-fusion zone inter-facial region was prepared using a focusedion beam (FIB) and in-situ lift out tech-nique. After attaching the TEM foil to a

copper grid, final thinning was performedon the sample at an acceleration voltage of30 kV, followed by 10 kV, and 1 kV for thefinal polishing step to get a 100-nm-thickTEM sample. The TEM studies were per-formed with a JEOL 2010F TEM

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Fig. 2 — A — Schematic of the laser brazing system used for joining AZ31Mg and Ni electro-plated steel sheets in the single-flare bevel lap joint config-uration showing the position of two thermocouples used for temperaturemeasurements; B — schematic of the 5-mm-wide tensile shear test specimen.

Table 3 — Composition of Ni Electroplating Solution and Electroplating Parameters

Plating Solution Composition (g/L) Electrodeposition Parameters

NiSO4•6H2O 263 Cathode current density 45–120 mA/cm2

Na2SO4 215 Time 5–20 minH3BO3 31 pH 3

Temperature 25°CAnode Graphite (8 cm2)Cathode Carbon Steel (6 cm2)

Fig. 3 — A laser-brazed Ni electroplated steel/AZ31B joint made using 8mm/s travel speed and 2.2-kW laser beam power: A — Top bead; B —transverse section of the joint.

A

A

B

B


equipped with an EDS.As shown in Fig. 2B, 5-mm-wide rec-

tangular-shaped specimens were cut fromthe brazed joints and subjected to tensile-shear tests with a crosshead speed of 1mm/min. Shims were used at each end ofthe specimens to ensure shear loads in thelap joint while minimizing induced cou-ples or bending of the specimens. Averagetensile shear strength was calculated fromtensile specimens to estimate the static

mechanical resistance and joint efficien-cies of the joints.

Results

A photograph of a laser-brazed Ni elec-troplated steel/AZ31B joint and a typicalcross-sectional view of the joint are shownin Fig. 3. This brazed joint was made using2.2 kW laser power, 8 mm/s travel speed,and 0.2 mm beam offset to the steel side.

The joint exhibited a uniform brazed areawith good wetting of both base materials.Partial melting of the AZ31B base metalwas observed. In contrast, when bare steelwas used, no bonding occurred betweenthe steel sheet and the braze alloy fusionzone (FZ) and wetting of the steel by thebraze metal was very poor (Ref. 15). The5.5-μm-thick Ni electrodeposited layer onthe surface of the steel significantly im-proved the wetting of the steel by molten


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Fig. 4 — Transverse sections of a laser brazed joint. A — Optical micrographof the entire joint and SEM images in different positions along the steel-FZinterface; B — position A; C — position C; D — position E; E — position F.

Fig. 5 — Typical temperature vs. time profiles measured during laserbrazing at the top and bottom side of the joint.

A

C

E

D

B


Mg-Al filler metal. Detailed microstruc-tural analysis of the fusion zone andAZ31B Mg alloy after the laser brazingprocess has been reported in our previousinvestigation (Ref. 15). This paper focuseson microstructural analysis of the steel-fu-sion zone interface.

Microstructural Evolution along the Steel-FZ Interface

Figure 4 shows the microstructure atdifferent locations of the steel-FZ inter-face. The Ni coating was not detected as aseparate layer along the interface after theLBP, which would suggest that it had en-tirely melted and gone into solution in theliquid immediately adjacent to the inter-face. It was observed that the microstruc-ture of the FZ-steel interface changed sig-nificantly across the FZ-steel interfacefrom the bottom (Position A, Fig. 4B) totop (Position F, Fig. 4E) side of the joint.In order to explain this change of mi-crostructure during the laser brazingprocess, temperature distribution acrossthe interface vs. time was measured duringlaser brazing using two thermocouples,one attached to the top side and the otherto the bottom side of the steel sheet (seeFig. 2A). According to the measured tem-perature profiles shown in Fig. 5, the steelsheet experienced maximum tempera-tures of 1151.1° and 652.7°C on the topand the bottom side, respectively. There-fore, a 500°C temperature gradient wasmeasured between the top and bottomside of the steel sheet during the laserbrazing process, since the laser beam wasfocused on the top of the filler metal, asshown in Fig. 2A (Ref. 15). This tempera-ture difference and gradient across thejoint interface during the laser brazingprocess is believed to be the main reasonfor the prominent change of microstruc-ture across the FZ-steel interface.

As shown in Fig. 4B, at the bottom ofthe interface a few diamond-shaped brightphases were formed near the steel-FZ in-terface. In order to identify these phases,a TEM foil was prepared from position Bof Fig. 4A. Figure 6 shows the TEM im-ages, EDS plot, and selected area diffrac-tion pattern (SADP) of these submicronparticles. The diffraction pattern shows astandard diffraction pattern of AlNi (withBCC structure) with [011] zone axis of theparticle. According to an EDS analysis ofthe diamond-shaped bright phases shownin Fig. 4B, the composition of the particleswas 49.6 ± 1.3 at.-% Ni, 45.4 ± 4.7 at.-%Al, and 5.0 ± 2.5 at.-% Mg, thus confirm-ing that the diamond-shaped particleswere mainly composed of AlNi inter-metallic compound (IMC). Representa-tive concentration profiles of Ni, Al, andMg across one AlNi particle are shown inFig. 6D, which indicates that a trace

amount of magnesium was found in thisparticle. It has been reported that each ofthe Al-Ni binary intermetallics has somesolubility for substitutional magnesiumatoms (Ref. 19).

Figure 7 shows the XRD spectra ob-tained from the middle of the steel-FZ in-terface. The area covered by the X-raybeam was a 300-μm-diameter circle. ThisXRD result confirmed the existence ofAlNi IMC, Fe, β-Mg17Al12, and α-Mg.The AlNi IMC compound was not foundat the middle of the FZ area, whereas theXRD pattern in Fig. 7 showed some weakpeaks suggesting that AlNi IMC hadformed mainly at the steel-FZ interface.

It was observed that upon moving fromthe bottom to the middle of the interface,which was associated with increasing tem-perature, the morphology of the IMC phasealong the interface changed from the dia-mond-shaped AlNi to a faceted dendritic-shaped phase (see Fig. 4C, D).

Energy-dispersive X-ray spectrometeranalysis results indicated this dendriticphase contained 43.0 ± 1.6 at.-% Ni, 52.1 ±2.0 at.-% Al, and 4.9 ± 0.5 at.-% Mg. Thiscomposition again corresponded with theAlNi IMC phase. In this area, the first pre-cipitated phase from the liquid was AlNiIMC, the same as at the bottom of the joint.This phase grew steadily in a faceted den-

dritic shape. As the interface temperatureincreased with moving from position A toposition E in Fig. 4A, the growth morphol-ogy of the AlNi phase changed from dia-mond-shaped to a faceted dendritic shape,as demonstrated in Fig. 4D. Continuousgrowth of the AlNi was observed in this areawith some dendrites having long secondarydendrite arms (see Fig. 4D).

At the top of the joint (position F in

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Fig. 6 — AlNi particle characterization at position B shown in Fig. 4A: A, B — TEM images; C — SADP inthe [011] zone axis of this particle; D — EDS composition line scans across an AlNi particle indicating linescans of Ni, Al, and Mg.

A

C D

B

Fig. 7 — X-ray diffraction pattern of the steel-FZinterface.


Fig. 4A), the morphology of the interfacialphase changed further and a high volumefraction of a particle-like phase with thecomposition of 48.4 ± 1.4 at.-% Ni, 50.1 ±1.2 at.-% Al, and 1.5 ± 0.4 at.-% Mg wasdetected (Fig. 4E). This phase was alsofound to be AlNi IMC phase. It shouldalso be noted that formation of the AlNiphase consumed almost all of the Al con-tent of the melt near the steel-FZ inter-face. Thus, no β-Mg17Al12 was observednear the interface compared with the cen-tral part of the FZ.

Solidification of the Remaining Meltbetween the AlNi IMC Phase and Steel

At the bottom of the joint, AlNi IMCfirst crystallized from the liquid close tothe interface and then supersaturated α-Mg solid solution containing 10.6 ± 3.6

at.-% Ni, 3.2 ± 1.7 at.-% Al, and 2.9 ±1.5 at.-% Fe (dark regions in Figs. 4B and8) solidified from the liquid during coolingalong with the AlNi phase. In some loca-tions between the bottom and middle ofthe interface, some gray lamellar phases,as shown in Figs. 8C and 9, were also ob-served between the AlNi IMC layer andthe steel. Figure 9A shows the position ofAlNi precipitates with respect to the steel-fusion zone interface in the prepared sam-ple during focused ion beaming for TEManalysis. Figure 9B, C shows TEM imagesof this lamellar (plate-like) phase. Ac-cording to EDS analysis, the white lamel-lae corresponded to α-Mg and the darklamellae containing 27.6 ± 7.2 at.-% Niand 72.3 ± 7.3 at.-% Mg represented theMg2Ni stoichiometric intermetallic com-pound (also confirmed by SADP analysis).Based on these results, these two phases

next to each other are the Mg-Mg2Nilamellar eutectic.

Formation of the Mg-Mg2Ni lamellareutectic was not uniform and continuousalong the interface. As shown in Fig. 8A,B, in some locations between the steeland the AlNi IMC layer, Mg2Ni crystal-lized in the form of a lamellar gray phaseand in other locations it was not seen anda dark solid solution of Mg containingsmall amounts of Ni, Al, and Fe wasformed. In the middle portion of the in-terface, the AlNi phase crystallized firstin the liquid (Fig. 4D). Then, dark α-Mgsolid solution containing 5.8 ± 2.1 at.-%Ni, 1.2 ± 0.3 at.-% Al, and 3.1 ± 0.5 at.-% Fe formed during cooling along withAlNi phase. Finally, at the top of thejoint, the AlNi phase precipitated heavilyin the liquid along the interface and thenthe remaining liquid solidified duringcooling in the form of α-Mg solid solution(containing 2.4 ± 0.6 at.-% Ni, 0.3 ± 0.1at.-% Al, and 3.4 ± 0.3 at.-% Fe) alongwith and among AlNi particles (see Fig.4E). Upon moving from the bottom tothe top of the interface, the Fe content ofthe remaining liquid between AlNi IMCand the steel increased from 2.9 to 3.4 at.-% due to more diffusion of Fe from thesteel side to the FZ at higher tempera-ture. In contrast, Al and Ni showed op-posite behaviors due to an increase in thethickness of AlNi IMC from the bottomto the top portion of the joint (from 5 to30 μm).

Transition Layer

Based on the TEM analysis, the AlNiphase did not grow epitaxially on the steelsubstrate, but instead nucleated and grew inthe liquid adjacent to the interface and wassurrounded by either α-Mg + Mg2Ni eu-tectic phases or just α-Mg phase close to theinterface. On the other hand, while it mayappear in Fig. 8 that all of the electroplatedNi had melted and gone into solution in theliquid and the α-Mg may have nucleatedand grown epitaxially from the steel surface,it is well known that Mg and Fe are an im-miscible couple. From a crystallographicpoint of view, it is not possible for magne-sium to nucleate on steel due to the verylarge lattice mismatching of Fe and Mg(Ref. 4). Therefore, another layer or phasemust be responsible for bonding betweenthe steel and fusion zone.

Further high-magnification mi-crostructural analysis of the steel-fusionzone interface was performed by TEM tofind an explanation for the observed in-terfacial phases. Figure 10A shows a TEMimage of the steel-fusion zone interface. Acontinuous nano-interlayer (50–200 nmthick) phase was observed along the inter-face, which was bonded to the steel side onone side and to the fusion zone on theother side. Higher magnification of this


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CH Fig. 8 — SEM images of the steel-FZ interface show the

solidification morphology of remaining melt betweenthe IMC layer and steel side: A — Position A in Fig. 4Anear bottom side; B — position B in Fig. 4A; C — Mg-Mg2Ni eutectic phases.

A

C

B

A

B

C

Fig. 9 — A — TEM sample attached to a coppergrid; B, C — TEM images of the lamellar phasesformed along the steel-FZ interface.


layer (as shown in Fig. 10B) confirmedgood coherency between this layer andsteel as well as the fusion zone. Accordingto EDS point analysis, the Ni content ofthe transition layer varied between 17 and40 at.-%. Figure 10C shows the selectedarea diffraction pattern (SADP) on thetransition layer that identified it as Fe(Ni)solid solution with face-centered cubic(FCC) structure. Therefore, this layerproved to be the key factor for realizing ametallurgical bond between the steel andfusion zone. Representative concentra-tion profiles of Fe, Ni, and Mg across theinterface between the fusion zone and thesteel are shown in Fig. 10D. It is evidentfrom these line scans that Fe, Ni, and Mgdiffused into each other as a result of thehigh temperature experienced during thelaser brazing process. As a result, two dif-fusion or transition layers formed betweenthe steel and fusion zone. According to theelement distributions of Fe, Ni, and Mg(see Fig. 10D), in transition layer I with athickness of almost 70 nm from the steelside, the Fe content decreased graduallywhile the Ni content increased. In thislayer, solid-state diffusion of Ni and Feinto each other is believed to control theoverall thickness of this layer.

Another diffusion layer (transitionlayer II) was observed in Fig. 10D betweenthe transition layer I and the fusion zone.The thickness of this layer was ≈ 60 nm. Aslight diffusion of magnesium from fusionzone into transition layer II was detected.It would appear that there was sufficientsolubility of the Mg in this Fe(Ni) inter-layer for diffusion of the Mg to occur andthat wetting and bonding of the α-Mg +Mg2Ni eutectic phases had in fact oc-curred with the thin Fe(Ni) interlayer thathad formed during laser brazing, and notdirectly with the steel.

Mechanical Properties

Due to the nonsymmetric configura-tion of the 5-mm-wide tensile-shear testspecimens (see Fig. 2B), a combination ofshear and tensile forces existed at the in-terface. Consequently, the joint strengthsare reported here as fracture load, since itis not possible to separate tensile andshear stresses. The average tensile shearstrength of the laser brazed steel-Ni-AZ31B joints using the Mg-Al filler metalwas found to be 153.7 ± 2.7 kgf (or 1506.3± 24.5 N). This is 153% higher than ten-sile shear strength of the laser brazed Al-coated steel-AZ31B Mg alloy specimensobtained in our previous study (Ref. 15).The low standard deviation of the tensileshear strength of the laser brazed steel-Ni-AZ31B joints (±2.7 kgf) compared withthe laser brazed steel-Al-AZ31B joints(±11 kgf) indicated that the laser brazingprocess was inherently stable and repro-

ducible. If only the shear plane is consid-ered, the average shear strength of thejoints was 96.8 MPa, or 60% of that ofAZ31B-H24 Mg alloy base metal.

All tensile-shear specimens fracturedin the FZ very close to the steel-FZ inter-face. Typical fracture surfaces of both thefusion zone side and steel side after tensileshear testing are shown in Fig. 11. Figure11A, C are low-magnification SEM micro-graphs of the fracture surfaces of the fu-sion zone side and steel side, respectively,and dimples are shown at high magnifica-tion in Fig. 11B, D. This uniform distribu-tion of the dimples is characteristic of duc-tile fracture surfaces. These fracturesurfaces indicated that the specimensfailed under conditions similar to tensiletest with a strong shear stress component(tensile-shear test). The effect of shearstress on the morphology of the dimples isvery evident in these micrographs. Thevertical direction in each of the micro-graphs is parallel to the direction of theshear, and the elongation of the dimplesunder the action of shear stress is evidentin Fig. 11B, D. The AlNi IMC compoundwas not found at the fracture surfaces.

The EDS analysis results of the frac-ture surfaces of both the steel and FZ side

also indicated that crack propagation dur-ing the tensile shear tests had occurred en-tirely in the FZ. Based on the EDS results,the composition of the fracture surface forboth steel side and FZ side were similar tothe FZ, meaning fracture passed throughthe FZ near the steel-FZ interface.

Figure 12 shows the XRD pattern fromthe fractured surface of the joint on thesteel side. Fe, α-Mg, and AlNi peaks wereseen in this X-ray diffraction result. Thesefindings were consistent with the SEM andEDS analysis results.

Discussion

From the above results, the interactionbetween the filler metal and surface of theNi-plated steel can be explained as follows(see Fig. 13):

Firstly, the solid-state Ni-plated steel isin contact with the liquid filler metal (Mg-Al alloy) at the laser brazing temperatureand, subsequently, the liquid Mg-Al alloyflows over the Ni surface — Fig. 13A.

Secondly, dissolution and diffusion ofNi atoms into the liquid occur, as shown inFig. 13B. At the same time, some solid-state diffusion of Ni atoms into the steelalso occurs. A slight diffusion of Fe atoms

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A

C

B

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Fig. 10 — A— TEM image of the steel-fusion zone interface; B— higher magnification of the selectedsquare area in A; C — SADP in the [011] zone axis of the interfacial phase; D — EDS line scan analy-sis of Fe, Ni, and Mg at the steel-fusion zone interface.


into the liquid was also observed. Mean-while, Mg atoms from the liquid mayslightly diffuse into the Ni-alloyed steel

side. Therefore, a thindiffusion or transitionlayer forms continu-ously along the inter-face between steel andfusion zone from thebottom side to the topside of the joint. Thistransition layer is asolid solution of Ni inFe (with low content ofMg for transition layerII). This Fe(Ni) solidsolution with FCCcrystal structure ismore favorable forbonding to Mg thanhaving a body-cen-tered cubic (BCC)phase along the inter-face. Zhang et al. (Ref.20) used an edge-to-edge matching crystal-lographic model topredict all orientation

relationships between crystals that havesimple hexagonal close-packed (HCP)and BCC structures, and they found that

the lattice mismatching of HCP (Mg) andBCC (Fe) is very large. On the basis of theobservation in our study, a diffusion layercomposed of Fe and Ni with FCC struc-ture can provide the conditions for theheterogeneous nucleation of α-Mg duringsolidification. The result is formation of ametallurgical bond between the steel andmagnesium alloy. A recent study by Liu etal. (Ref. 4) showed that a nano-layer ofFe2Al5 on steel can also be a transitionlayer to bond Fe to Mg due to the low en-ergy interfaces and good match of latticesites between Fe and Fe2Al5 as well as Mgand Fe2Al5. The same behavior was ob-served for the Fe(Ni) transition layer inthis study. Formation of a transition zonewas also reported in other studies (Refs.5–8) using different joining techniques,when an interlayer was used between steeland Mg alloy. These transition layers onsteels were reported to make it possible tojoin Mg and steel.

Thirdly, during the solidificationprocess, the AlNi phase with a high melt-ing point (1133°C) precipitates from theliquid and grows in a form of faceted den-drites very close to the interface — Fig.


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Fig. 11 — SEM images of typical fracture surfaces after the tensile shear test. A, B — Fusion zone side at different magnifications; C, D — steel side at differ-ent magnifications.

Fig. 12 — X-ray diffraction pattern of the fracture surface of the steel side.

DC

BA


13C. These faceted dendrites form due tokinetic difficulties in forming new planesof atoms (Ref. 21). In this type of dendrite,the growing direction of dendrite arms areones that are capped by relatively slowgrowing planes (usually low-index planes)(Ref. 21). The slowest growing planewould be expected to be the closest-packed planes. Weinberg and Chalmers(Ref. 22) reported that the axis of a pyra-mid, whose sides are the most closelypacked planes, is generally the major den-drite direction. As a result, for AlNifaceted dendrites with BCC structure, thisdirection is <100>. Therefore, theprocess of solidification at the middle partof the joint starts with the nucleation andgrowth of the AlNi-faceted dendritesalong the <100> growth direction.

Fourthly, if the Ni content of the re-maining liquid between the steel side andformed AlNi precipitates is high enough,Mg2Ni with a melting point of 762°C nu-cleates (see Fig. 13C). Formation ofMg2Ni depends on sufficient Ni concen-tration in the remaining liquid near thesteel-FZ interface after formation of theAlNi IMC. The Ni content of the remain-ing liquid after precipitation of AlNi in-creases from 2.4 at.-% at the top side ofthe interface to 10.6 at.-% at the bottomportion because formation of the AlNiIMC layer consumed the Ni atoms nearthe interface and the volume fraction ofthis phase increased from the bottom tothe top portion of the joint.

Based on the above analysis, highenough concentration of Ni in the re-maining liquid close to the bottom side ofthe joint after formation of AlNi IMC re-sulted in formation of the Mg2Ni + α-Mglamellar eutectic in the form of a grayphase between the AlNi IMC and steel. Inorder for this lamellar eutectic to grow,the local composition of the fusion zoneshould be close to the eutectic composi-tion (10 at.-% Ni, according to the Mg-Nibinary phase diagram) (Ref. 21). Reac-tions between Mg in the fusion zone andNi along the interface caused formation ofthe Mg-Ni eutectic phase. This reactioncan be represented by the following bal-anced chemical reaction:

L(10 at.-% Ni)←508°C→Mg2Ni(33 at.-% Ni)+Mg(0 at.-% Ni) (1)

Therefore, at the bottom of the inter-face, two reactions occurred; the first onewas precipitation of AlNi from the liquidand the second was the eutectic reactionbetween Mg and Ni in the FZ (reaction 1).In the case of reaction sequences, firstAlNi forms near the interface and then theremaining liquid with a low Al content be-tween the AlNi IMC and steel-FZ inter-face, which is still rich in Ni, undergoes aeutectic reaction with Mg and results in

the formation of the lamellarα-Mg + Mg2Ni eutectic.

With the formation of theAlNi IMC layer, diffusion ofNi atoms from the steel side tothe FZ is blocked. Therefore,the concentration of Ni in theremaining liquid between theinterface and preformed AlNiphase is expected to be higherthan the remaining liquid onthe other side of the AlNiphase. The result is the forma-tion of Mg2Ni just betweenthe AlNi phase and steel (seeFig. 8A, B).

In the top portion of theinterface, with the nucle-ation and growth of the AlNiparticles, most of the Niatoms are consumed. There-fore, the Ni content of theremaining liquid would notbe enough for formation ofthe Mg2Ni phase.

Conclusions

1. With the addition of anelectrodeposited Ni inter-layer on steel sheet, singleflare bevel lap joints ofAZ31B-H24 Mg alloy tosteel sheet were renderedpossible by the laser brazingprocess, and a uniformbrazed area with good wet-ting and bonding of both basemetals was achieved.

2. Dissolution of the Nicoating layer during the laserbrazing process led to theformation of new AlNi IMCphases and also a Mg-Ni eu-tectic zone along the inter-face. The AlNi intermetalliclayers at the steel-FZ inter-face formed in the sequenceof diamond-shaped, den-dritic, and nodules from thebottom to the top portion ofthe joint.

3. The formation of anano-scale Fe(Ni) transitionlayer on the steel by solid-state interdiffusion betweenFe and Ni during laser braz-ing was found to be responsi-ble for the formation of ametallurgical bond betweenthe steel and the Mg-Al braz-ing alloy.

4. The average shear strength of thejoints reached 96.8 MPa, 60% that of thebase metal of AZ31B Mg alloy. Fracturesurface analysis showed that fracture oc-curred in the FZ close to the steel-FZ interface.

Acknowledgments

The authors wish to acknowledge sup-port of the American Welding Society(AWS) Graduate Fellowship program, theNatural Sciences and Engineering Re-search Council of Canada (NSERC), and

9-sWELDING JOURNAL

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Fig. 13 — Formation of transitional layer and intermetallic com-pounds during laser brazing of Ni-plated steel-AZ31B with Mg-Alfiller metal: A — Wetting of the Ni-plated steel by molten filler metaland dissolution and diffusion of Ni into the FZ and steel substrate; B— formation of the transitional layer and aggregation of Ni along theinterface; C — nucleation and growth of AlNi IMC, and epitaxialgrowth of the remaining liquid in the form of α-Mg + Mg2Ni eutecticonto the thin Fe(Ni) interlayer.



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Magnesium Network of Canada (Mag-NET) for sponsoring this work. The au-thors would like to acknowledge the help-ful comments of Dr. Scott Lawson fromthe Centre for Advanced Materials Join-ing at the University of Waterloo.

References

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2. Mao, H. K., and Bell, P. M. 1979. Equa-tions of state of MgO and Fe under static pres-sure conditions. Journal of Geophysical Re-search 4533–4536.

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5. Zhao, X., Song, G., and Liu, L. 2006. Mi-crostructure of dissimilar metal joint with mag-nesium alloy AZ31B and steel 304 for laser-tungsten inert gas lap welding. Transactions ofChina Welding Institution 27: 1253–1256.

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AWS Expands International Services

With international membership on the rise, the American Welding Society (AWS)launched a series of country-specific Web sites known as microsites for members toaccess information in their native languages.

Multilingual microsites are now live for Mexico at www.aws.org/mexico, China atwww.aws.org/china, and Canada (English/French) at www.aws.org/canada. They fea-ture information on services offered by AWS in each country, membership benefits,exposition information, online education, and access to AWS publications and tech-nical standards.

Other countries will be added later.


welding journal | january 2013

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