WELDING RESEARCH Development of Requirements for...

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ABSTRACT. The resistance spot weldingof two galvanized DP600 steels with 1.8and 2.0-mm sheet thickness was investi-gated to enhance our understanding ofweld fracture during a well-established au-tomotive quality control test. The effectsof several process parameters (includingweld current, weld time, and weld force)on zinc expulsion, mechanical properties,defects, and microstructures of resistancespot welds were all studied in relation withthe two major types of weld fracture (i.e.,interfacial and button pullout). It was es-tablished that the slight differences inchemical composition and galvanizedcoating between the two selected DP600steels affected negligibly weld microstruc-tures, and were of no measurable conse-quences to weld fracture. In contrast, thedifference of 10% for their sheet thicknesswas responsible for smaller weld diame-ters, lower weld tensile-shear forces,larger shrinkage voids, and more frequentinterfacial fractures in welds from thethicker DP600 steel. For the two steels, oc-currence of weld interfacial fracture waseliminated using long weld times (>20 cy-cles), low currents (<9 kA), and highforces (>900 lb, or 4.0 kN); i.e., processparameters that increased weld diameterswhile preventing zinc ingestion into the fu-sion zone. The effect of zinc was mostprominent in welds that were made ab-normally fast (e.g., ~5 cycles), where bothsolidification cracking and a change intype of weld fracture were found due to in-

terdendritic low-melting point con-stituents.

Introduction

The drive for vehicle mass reductionhas led to the introduction of new engi-neering materials. The Ultra Light SteelAuto Body (ULSAB) project has shownthat car body mass can be reduced by 25%using advanced high-strength steels(AHSS) and innovative processes (Ref. 1).Among the AHSS, the dual-phase (DP)steels have been the subject of particularattention owing to their good combinationof high strength and ductility. The termdual-phase steel refers to the predomi-nance of two phases, the body-centered-cubic (bcc) a-ferrite and the relativelyharder body-centered-tetragonal (bct)martensite. These two phases are pro-duced by some annealing in the A1–A3“intercritical” temperature, where austen-ite and ferrite are formed, and a subse-quent rapid cooling where the austenite iseventually transformed into martensite.Since suitable amounts of ferrite andmartensite in DP steels may be producedfrom a combination of heat-treating para-meters, compositions of DP steels mayvary significantly between steel makers, as

found in Part 2 of this study (Ref. 4) to bepublished in the near future. To createyield strengths of 350 MPa and ultimatestrengths of 600 MPa, the DP600 steels (ascommonly referred by the ULSAB Con-sortium.) (Ref. 1) were all found to pos-sess about 15% martensite. Compared toprecipitation-strengthened or solid-solution-strengthened HSLA steels,DP600 steels exhibit a slightly lower yieldstrength, a continuous yield behavior dueto enough active slip systems in the ferritephase, and a more uniform and highertotal elongation (over 21%) (Refs. 1, 2).These last two properties explain their rel-atively good formability; a property, whichcombined with their high strengths, hasmade DP steels attractive for automotiveapplications (Refs. 1–3).

The acceptance of DP steels in manu-facturing environments has been ex-tremely gradual though DP steels offer ev-ident advantages, and they havedemonstrated a good weldability (Refs. 3,5–8). However, on automotive factoryfloors and in testing laboratories, the def-inition for weldability can differ, and soare the tests and criteria used to differen-tiate “good” welds from “bad” welds. Inthe simplest of all quality-control tests, ahandheld chisel is inserted in betweenspot-welded coupons to force the welds tofracture, and, based upon a visual inspec-tion, determine if the same welds would beappropriate in vehicle applications (Ref.9). In early chisel tests, spot welds inDP600 steels had been found to under-perform welds in traditional automotivesteels; i.e., non-AHSS (Ref. 1) (e.g., inter-facial-free steels, low-carbon steels,rephosphorized steels, mild steels, orHSLA steels). Especially with the thickerDP600 steel gauges, welds were seen tofracture in the same plane as the sheet sur-

Development of Requirements forResistance Spot Welding Dual-Phase

(DP600) SteelsPart 1 — The Causes of

Interfacial FractureWeld fracture was investigated in relation to weld

parameters and steel sheet characteristics

BY M. MARYA and X. Q. GAYDEN

KEYWORDS

Automotive SteelsDual-Phase (DP600)Resistance Spot WeldingPropertiesMicrostructuresDefects

M. MARYA, formerly with G. S. Ansell Depart-ment of Metallurgical and Materials Engineering,Colorado School of Mines, Golden Colo., is nowwith NanoCoolers, Inc., Austin, Tex. X. Q. GAY-DEN is with Materials and Process Laboratory,General Motors Research and Development, War-ren, Mich.

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faces, causing the so-called “interfacialfracture.” The occurrence of weld interfa-cial fractures in DP600 steels was in con-trast with the traditional automotivesteels, where weld fracture is known tooccur consistently in the fusion zone pe-riphery, promoting the so-called “weld

button.” For spot welds to be reliable dur-ing vehicle lifetime, they are required totear a button during quality control. Theformation of a weld button during qualitycontrol indeed indicates that the sameweld would have been able to transmit ahigh level of force, thus cause severe plas-

tic deformation in its adjacent compo-nents, and increased strain energy dissipa-tion in crash conditions.

In Part I of this investigation on theweldability of DP600 steels in productionenvironments, fracture in spot welds fromtwo commercial steels was induced mainly

Fig. 1 — Secondary electron images showing the galvanized layers on thetwo DP600 steels and electron dispersive spectra indicating their approxi-mate chemical compositions.

Fig. 2 — Process map describing the effects of the current and weld time onzinc expulsion for both DP18 and DP20 (1200-lb weld force).

Fig. 3 — Process map describing the effects of the weld force on the zincexpulsion for the DP20 steel (1200-lb weld force).

Fig. 4 — Reduction in thickness or thinning vs. current and weld time for weldsmade in DP18 and DP20 steels (1200-lb weld force).

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by the chisel test (Ref. 9). When this studywas initiated in 2003, it was primary to firstestablish a relationship between processparameters and occurrence of weld inter-facial fracture using a minimum of twoDP600 steels. Since weld diameter is well-known to be a key player in weld fracture(Refs. 6, 9–14), weld diameters were mea-sured and correlated to the process para-meters, as were the weld microstructures.In this first study, the weld microstructureswere thoroughly examined due to well-

established relationships with the me-chanical properties, and thus fracture(Refs. 5–8). In this article, the reader willdiscover that the weld microstructures arefirst discussed and linked to the steel com-positions, as suggested by the NipponSteel carbon equivalent (CE) here pre-sented as Equation 1 (Refs. 12–14). ThisCE, developed from non-AHSS spotwelds that were quasi-statically cross-tension tested (Refs. 12–14), connectschemical composition with one of the two

types of weld fracture. This CE, yet un-proven for AHSS steels, presented ameans for measuring DP steel spot weld-ability with a load orientation that was nottoo different from the chisel test load ori-entation. Equation 1 for this CE infersthat a relatively hard and brittle weld (i.e.,one that is made of martensite) havingenough sulfur and phosphorus (i.e., ele-ments that segregate to boundaries and in-crease solidification cracking susceptibil-ity, Refs. 15, 16) will fail interfacially

Fig. 5 — Weld diameter (as measured after peeling) vs. current for weldsmade in DP18 and DP20 steels (20-cycle weld time and 1200-lb weld force).

Fig. 6 — Fracture force (peak force recorded during tensile-shear testing)vs. current for spot welds made in DP18 and DP20 steels (20-cycle weldtime and 1200-lb weld force).

Fig. 7 — Process maps showing effects of both current and weld time on thetype of weld fracture encountered during chisel testing on DP18 (left) andDP20 (right) steels (1200-lb weld force).

Fig. 8 — Percent weld interfacial fracture vs. current and force for spot weldsmade on DP20 (20-cycle weld time).


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during testing if CE is greater than 0.24(Refs. 12–14).

(1)

Although not detailed here, but sug-gested elsewhere (Ref. 14), zinc contami-nation from the coatings is also anotherpotential contributor to weld fracture.Consequently, the effects of process para-meters, sheet thickness, steel composi-tion, and coating were all important toconsider for fully explaining the weldabil-ity of DP600 steels, since they all influenceweld dimensions, weld microstructures,and weld mechanical properties, includingfracture characteristics (Refs. 4–19).

ExperimentalProcedure

Materials

Following suggestions by manufactur-ing engineers, two hot-dipped galvanizedDP600 steels of commercial availabilitywere selected with two distinct gauges.The thinner steel, designated as DP18,was 1.8 mm thick, whereas the second,DP20, was supplied in a 2.0-mm sheetthickness. For testing, the steel sheetswere sheared into 38 ¥ 127-mm coupons.Guided by Equation 1, the steel chemicalcompositions were precisely measured

and verified using acombination of tech-niques: X ray fluores-cence (XRF), induc-tively coupled plasma(ICP) spectroscopy, andwavelength dispersivespectroscopy (WDS).

Welding

All resistance spotwelds were fabricatedusing a mid-frequencyDC welding machine,where current was de-livered through C15000Cu-Zr alloy caps with a5.0-mm face diameter.To minimize weld-to-weld variations intro-duced by cap wear, thecaps were regularly in-spected using referencewelds, which were gen-erated each time newcaps were inserted.These reference weldswere all produced usingtypical process parame-ters (e.g., 9 kA, 20 cyclesand 1200 lb, or 5.3 kN).When the repeat of a reference weld withused caps resulted in a weld with a new di-ameter and type of weld fracture, new capswere systematically mounted. This pre-caution guaranteed that the welding con-

ditions were not drifting over time, andthat the two steels could be comparedunder identical conditions at any time ifneeded.

Each welded sample was made of two

CE CSi30

Mn20

2P 4S0.24limit = + + + +

Fig. 9 — Microhardness profiles for welds in DP18 and DP20 producedusing widely different welding conditions (microhardness was measuredusing a 200-g load and 30-mm away from the weld centerline).

Fig. 11 — X-ray compositional maps near the center of an unetchedspot weld made in DP20 steel (12.5 kA current and 20-cycle weld time)with line scans for three alloying elements, all revealing some mi-crosegrgation.

Fig. 10 — Optical micrographs of microstructures in a spot weldmade in DP20 steel (12-kA current and 20-cycle weld time), and EDSresults revealing the presence of nontransition metals, including sulfur.

A B

C D

E


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identical coupons. On each, three spotwelds, separated by about 35 mm, wereproduced. Weld current and weld timewere varied between 7 and 40 kA and be-tween 2 and 40 cycles, respectively. Theweld force (i.e., the compressive force in-duced by the electrodes) was applied 90cycles before the current, and lasted 5 cy-cles longer. Although several force levelswere selected, primarily to differentiatethe two DP600 steels, the weld force waspreferentially held at 1200 lb (5.3 kN).

Property Measurements

The overwhelming majority of thewelded samples was either peeled to forcethe formation of weld buttons and mea-sure their diameters, or chisel-tested to

identify the predominant type of fracture(Refs. 9, 10). Other welded samples weretensile-sheared at a low-crosshead veloc-ity of 4 mm/min to precisely capture andquantify weld static characteristics. Thistest was preferred over cross-tension andcoach-peel tests because of greater datareproducibility, reliability, and accuracy;all necessary to compare the two DP600steels. For tensile-shear testing, the weldswere centered in the overlap section,which like the coupon width was 38-mmlong. The same test configuration is de-scribed in other studies (Refs. 5, 6).

Several welds were cross-sectioned,cold-mounted, ground, and polished withcarbide paper, diamond paste, and finecolloidal silica. The microstructures werethen revealed by etching with a 2% nital

reagent and examined by optical and scan-ning electron microscopy (SEM). Chemi-cal compositions, including high-resolution mapping, were measured usingelectron/wavelength dispersive spec-troscopy (EDS/WDS). The Vickers mi-crohardness under a 200-g load was alsomeasured, and linked to the observed microstructures.

Results and Discussion

Base Material Characteristics

Measured chemical compositions,summarized in Table 1, show that theDP18 steel had slightly more carbon thanthe DP20 steel. This greater carbon contentof DP18 appeared to counterbalance a rel-

C D

A B

Fig. 12 — Secondary electron images of an interfacially failed spot weld in DP20 steel (12 kA current, 10-cycle weld time, 1200 lb weld force).

Table 1 — Measured Chemical Composition of the Two DP600 Steels (wt-%) with CE, as Defined by Equation 1

C Mn Cr Mo Ni Ti V Al Si P S B N O CE

DP18 0.13 1.40 0.21 0.05 0.02 0.01 <0.01 0.02 0.4 0.02 <0.002 <0.001 0.0051 0.0025 0.253

DP20 0.11 1.50 0.27 <0.005 0.02 <0.005 0.06 0.06 0.1 0.02 <0.003 <0.001 0.0065 0.0028 0.228


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ative shortfall of strengthening elementssuch as manganese, chromium, and vana-dium. In both DP18 and DP20, the boroncontents were also less than 10-wt-ppm; aconcentration that therefore neither couldincrease steel hardenability (Ref. 17) norcracking susceptibility (Ref. 14).

Table 1 provides the carbon equiva-lents for DP18 and DP20. Although theCE was determined with Equation 1 thatnormally applies to non-AHSS steels(Refs. 12–14), the CE could still be uti-lized as a relative indicator of hardenabil-ity and weldability. Since the CE for bothsteels was close to the limit of 0.24, vari-ability in the types of weld fracture was ex-pected to be greater than in non-AHSSsteels, where the CE is typically smaller.Also, note that the carbon equivalent ofDP18 not only exceeded that of DP20, butalso the value of 0.24 given by Equation 1(Refs. 12–14). Consequently, with only theCE to describe weld fracture, the spotwelds in DP18 could have been expectedto fracture interfacially more frequentlythan in DP20. The question as to whetherthis statement is correct or not is answeredlater in this paper.

As for the as-received microstructures(not shown), they were comparable inDP18 and DP20. Both steels had averagegrain sizes of about 10 µm and containedapproximately 15% martensite. However,hardness of the DP18 steel was greaterthan that of DP20, averaging 205 kg/mm2

vs. 185 kg/mm2 for the DP20 steel. Coinci-dentally or not, this difference in hardness

agrees with the CE values.SEM observations revealed that the

galvanized layer was 9 µm thick in DP20and nearly twice as thick in DP18 (Fig. 1),while EDS and XRF measurementsshowed that coating compositions weresimilar with at least 99 wt-% zinc and lessthan 1 wt-% aluminum. Therefore, thepossibility that coating compositions af-fected weld fracture in the two DP600steels could be eliminated.

Spot Weld Expulsion

In the resistance spot welding of galva-nized steels, the expulsion of zinc (i.e., theelement that melts and vaporizes first) isimportant to examine, as it limits the se-lection of the process parameters. Zinccan be expulsed both at the electrode-sheet interfaces, and at the interface be-tween the two sheets where the weld forms(Refs. 18–20). Due to water cooling withinthe electrodes, the temperature is nor-mally greater at the interface between thesheets and expulsion occurs there first(Ref. 18). Because this expulsion is moredifficult to detect when it first occurs, onlythe expulsion from the electrode-sheet in-terface is considered here.

Figure 2 describes the effects of weldcurrent and weld time on zinc expulsion.To determine with precision the thresholdof zinc expulsion for the two steels, eachdata point was generated from at leastthree welds. Judging from Fig. 2, the pos-sibility that the thicker zinc coating of

DP18 (Fig. 1) could have promoted moreexpulsion was ruled out since a singleboundary between the expulsion-free andthe expulsion regions was found for thetwo steels. Although both DP18 and DP20were indistinguishable here, Fig. 2 waslater found to be extremely useful for se-lecting appropriate process parameters,and to further compare DP18 and DP20.

In Fig. 2, the solid line demonstratesthat the current at the beginning of expul-sion is inversely related to the weld time.Consequently, to create large welds, cur-rent and weld time must be selected suchthat they are positioned close to thisboundary line, and, to prevent expulsion,they must be also on the left of this bound-ary line. Moreover, to enhance weldprocess repeatability, the best-controlledvariable (i.e., weld time) must be extendedas much as possible, whereas the current,the second variable of Fig. 2, must be min-imized to alleviate its relative contributionin the weld formation, and to prevent ex-tensive cap wear (Ref. 6). Unlike the weldtime, the current is self-regulated andvaries as contact resistance, material prop-erties, alignment, and cap wear graduallychange from one weld to another (Refs. 6,15, 18–20). For this study, where a highlevel of reproducibility between welds wasneeded to compare DP18 and DP20, ex-tended weld times (> 20 cycles) and cor-respondingly small currents (<9 kA) wereparticularly appropriate.

Figure 3 is a complementary figure de-scribing weld force effect on zinc expul-

Fig. 13 — Percent voids in weld fusion zones vs. weld time for spot welds madein DP18 and DP20 steels (12-kA current and 1200-lb weld force).

Fig. 14 — Percent voids in weld fusion zones vs. current and weld forcefor spot welds made in DP20 steel (20-cycle weld time).


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sion with weld times of 20 and 30 cycles.For currents of 12 kA and less, zinc expul-sion was prevented by reasonably low-weld forces. In contrast, the currents over12 kA were found to be impractical as therequired forces to prevent expulsion werethen too high for many industrial weldingmachines. Despite this new constraint, we

occasionally selected currents of 12 kAand higher, as justified later.

Weld Joint Properties

Spot Weld Morphology

In addition to preventing zinc expulsion,

the compressive force applied by the elec-trodes forges the fused region to a thinnergauge. Such thinning caused by indenta-tion can alter spot weld properties, asquantified in a recent study (Ref. 21), andwas therefore another factor guiding theselection of the process parameters. For aweld force of 1200 lb (5.3 kN), Fig. 4 de-

Fig. 15 — Secondary electron images of a spot weld produced using unusual process parameters (24-kA current, 5-cycle weld time, 1200-lb weld force). Weldbutton formation appeared to be promoted by zinc ingestion.

A B

C D

E F


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scribes current and weld time effects onthe thinning seen across more than 60 spotwelds from DP18 and DP20. Exactly likein Fig. 2, the results for both DP18 andDP20 were nearly identical. In Fig. 4, notethat we superimposed a dash line to showregions of expulsion and no expulsion (leftto the dash line). At the onset of expulsion,thinning never exceeded 30%, and wasless than 20% for most welds. With thislimited indentation, numerous studieshave shown that the indentation effect isnegligible compared to that of weld diam-eter (Refs. 15, 20, 21). Consequently, theeffect of process parameters on weld di-ameter (measured after peeling to pro-mote weld buttons) had to be examinedbefore a link with the types of weld frac-ture could be researched.

Figure 5 shows that weld (button) di-ameter increases with the current at a ratethat is gradually and monotonically de-creasing, as observed elsewhere (Refs. 20,22–24). When the current exceeded 11kA, the weld button diameters in bothsteels stabilized slightly above 6.0 mm. Atlower currents, the welds were noticeablylarger in DP18, particularly when buttondiameters were less than 5.5 mm. This re-sult revealed that welds in DP18 formedat lower heat inputs than in DP20, andthat could be reasonably well explained bythe fact that DP18 was 10% thinner thanDP20. With thickness emerging as themain explanation for the discrepant welddiameters in DP18 and DP20, we were ex-pecting a number of spot welds fromDP18 and DP20 to exhibit different typesof fracture, even with identical processparameters. The effects of process para-meters on weld mechanical properties, in-cluding types of weld fracture, were there-fore important to clarify, first underwell-controlled conditions, and thenusing the chisel test; i.e., as if welds werequality controlled on factory floors.

Tensile-Shear Test Results

Many welds, fabricated with identicalparameters as in Fig. 5, were tensile-sheared to quantify their load-carrying ca-pabilities, and, despite loading conditionsdifferent than in the chisel test, determineif one steel was also more prone to spotweld interfacial fracture than the other. InFig. 6, peak force recorded during testing(i.e., the weld fracture force) is repre-sented as a function of current. Figure 6shows that the weld fracture forces wereconsistently greater in DP18. Althoughthe trend lines for the two steels werealike, the trend line for DP18 was alsoshifted to the left of that for DP20; a con-firmation that the welds began to form atsmaller currents in DP18, and that the

greater fracture forces seen in the weldsof DP18 were primarily due to its thinnergauge.

Among the welds in Fig. 6, six out ofthe 38 (all in DP18) created buttons. Withseveral, the fracture forces were seen toreach a maximum at intermediate cur-rents (i.e., 10 to 10.5 kA). At these cur-rents, interfacial fracture was also seen tofully disappear. With a further increase incurrent, interfacial fracture continued tobe avoided, but the sheets were alsoalarmingly indented. As confirmed by Fig.4, with currents near 10 kA, thinningacross welds was about 20%. For the weldsproduced with greater currents, thinningnot only exceeded 30%, but the weldswere also observed to form buttons alongthe indentation. The possibility that in-dentations over 20 to 30% influenced weldfracture during tensile-shear testing, andespecially during chisel testing, was there-fore raised.

Quality (Chisel) Test Results

Figure 7 compares chisel test resultsfor spot welds also made with a 1200-lb(5.3 kN) force. The two types of fracture,interfacial (white squares) and button-pullout (black squares), are mapped as afunction of current and weld time. Sinceeach data point was generated by testing 3to 12 welds depending upon the observedrepeatability, results from more than 400welds are summarized in Fig. 7. Whileboth steels exhibit regions where only asingle type of fracture occurred, the DP20steel is characterized by having a regionwith the two types of weld fracture. Thisregion is bound by currents between about10 and 14 kA and weld times betweenabout 15 and 25 cycles. As currents and/orweld times were increased, the percentageof each type of fracture varied from zeroto one hundred, and vice versa. For DP20,interfacial fracture was also found to van-ish when expulsion started. Of these twoobservations, the first indicated that thechisel test was largely reproducible, andthe second inferred that zinc expulsionwas potentially related to weld fracture, ashypothesized previously.

To determine if zinc expulsion had in-fluenced the type of weld fracture, somewelds were produced by varying currentsand forces, which, as seen in Fig. 3, bothaffect expulsion remarkably well. Thewelds, made with selected forces of 600 lb(2.7 kN), 900 lb (4.0 kN), 1200 lb (5.3 kN),and 1700 lb (7.6 kN), were subsequentlychisel tested, and correlations betweentype of weld fracture and occurrence of ex-pulsion were searched. The results ofthese tests are summarized in Fig. 8, wherepercentage of a given type of weld fracture

is represented as a function of current andforce. For a given force, Fig. 8 shows thatthe occurrence of weld interfacial fracturegradually decreased with the current.More specifically, at 600 lb (2.7 kN), ex-pulsion of zinc started before the currentcould be raised high enough to fully elim-inate interfacial fracture. In contrast, at1700 lb (7.6 kN), zinc expulsion was de-tected after interfacial fracture had beenprevented. For the process conditions ofFig. 8, it became clear that zinc expulsionwas not primarily related to weld fracture,and the common boundary in Fig. 7 for theregions of mixed fracture and expulsioncould only be coincidental.

Figure 8 also points out that the highweld forces reduced the minimum currentto prevent interfacial fracture. Indeed,once weld force exceeded 900 lb (4.0 kN),weld buttons were produced at smallercurrents (10 kA vs. 12.5 kA at 600 lb, or 2.7kN). This result confirmed that the level ofapplied force affected weld formation,presumably by affecting contact area andcontact electrical resistance (Refs. 18–20,24, 25). Since Fig. 8 also shows that weldcurrent was negligibly affected by forcesover 900 lb (4.0 kN), we concluded inagreement with another study (Ref. 19)that the contact resistances were practi-cally unchanged once a certain force, orpressure, is exceeded.

In Fig. 8, also note that the minimumweld diameter to eliminate interfacialfracture was consistently in the vicinity of5.5 mm regardless of weld force and cur-rent. This 5.5-mm value, measured afterchisel testing, was confirmed by some ofthe test data presented earlier. For cur-rents of 9 and 10 kA, Fig. 7 showed thatweld interfacial fracture fully disappearedonce weld times exceeded 25 cycles inDP18 and about 30 to 35 cycles in DP20.With these process parameters, Fig. 5 thenrevealed that interfacial fracture wouldnot occur for weld button diameters over5.5 mm in DP18 and 6.0 mm in DP20.Based upon this analysis and the data ofFig. 8, we therefore confirm that someminimum weld diameters could be recom-mended to prevent weld interfacial frac-ture, as further investigated in Part II ofthis work (Ref. 4).

As complementary remark resultingfrom a comparison of Figs. 6, 7, and 8, notethat many of the tensile-sheared welds ofFig. 6 that fractured interfacially wouldhave produced buttons if instead theywould have been chisel tested. Thisdemonstrates that the chisel test is less se-lective than the quasi-static tensile-sheartest, and that any weld identified as a“good” weld by the tensile-shear testwould also be a “good” weld if tested inproduction using the chisel test.


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Weld Heterogeneity

To complete this study, metallurgicalanalyses were needed in view of variousopinions that came to us. We had to de-termine if differences other than sheetthickness explained the greater suscepti-bility of DP20 toward weld interfacial frac-ture. Since steel chemical compositionfirst appeared to be important (Equation1), variations in hardness (i.e., strength),microstructure, and local compositionwere examined to attempt determining allcauses, primary or secondary, of spot weldfracture in DP600 steels.

Variations in Microhardness

For both DP18 and DP20, Fig. 9 de-picts microhardness profiles for weldsproduced with both regular and unusuallyfast welding schedules (12.5 kA, 20 cyclesvs. 40 kA, 2 cycles). The average hardnessvalues of the weld fusion zones are alsoshown by the horizontal straight lines. Av-erage hardness was higher in the rapidly-made high-current weld, as seen by com-paring the upper and lower graphs. Also,for all welds, microhardness fluctuatedslightly and always peaked in the HAZ,marked as HAZ. Near the fusion zone,where microstructure was the hardest,hardness exceeded 400 kg/mm2 and wasabout twice greater than in the base mate-rial. Like many non-AHSS, which have allconsistently produced buttons after chiseltesting, no measurable softening was de-tected near the base metal (Refs. 6, 27). Asoftening in the HAZ would have helpedplastic strain to localize and fracture tofollow outside the hard fusion zone (Refs.27–29). Therefore, this feature alonewould have explained why welds in DP600fracture more interfacially than tradi-tional automotive steels.

Regardless of the welding parameters,the HAZ and the fusion zone were harderin DP18, a result that agrees with the CEvalues of Table 1. While microhardness inboth fusion zone and HAZ were greater inDP18, the hardening with respect to theinitial microstructure (or percent increasein hardness) was greater in the welds ofDP20. This hardening was 2.17 in DP20 vs.2.05 in DP18. If comparing strength of fu-sion zone or HAZ with that of base metalis important, as suggested by mathemati-cal modeling (Refs. 28, 29), then examin-ing hardening was not helpful here in ex-plaining the greater susceptibility of DP20to weld interfacial fracture, because agreater hardening generally correlateswith less interfacial fractures (Ref. 27). If,instead, the CE of Equation 1 is used todifferentiate steels by types of spot weldfracture, then CE cannot explain the su-perior performance of DP18 during chisel

testing. Of all factors, only sheet thicknesshas to this point explained the dissimilarbehavior and different weldability ofDP18 and DP20.

Variations in Microstructure

Examination by optical microscopy ofweld cross sections neither distinguishedDP18 from DP20, but did reveal mi-crostructural features that potentially af-fected weld fracture in DP18 and DP20.Fig. 10 depicts four optical micrographs ofa characteristic spot weld in DP20. Thelow-magnification micrograph of Figure10A shows that microstructure was moreheterogeneous in the fusion zone than inthe HAZ. Figure 10B, for the colder sec-tion of the HAZ, demonstrates that its mi-croconstituents were considerably finerthan those of either the base metal or thefusion zone (Figs. 10C and 10D). This isexplained by the fact that austenitizingwas incomplete in the colder section of theHAZ, and even when austenite grainsformed, grain growth was restricted by thethermal cycles. In this section of the HAZwhere austenite grains are fine, the result-ing high density of grain boundaries con-stitute obstacles against the formation oflarge martensite laths. As opposed to thecoarser martensite of Fig. 10E, observedin the fusion zone, the fine martensite ofthe HAZ was not well resolved by SEM.As shown in Fig. 9, hardness correspond-ing to the fine microstructures of the HAZwas also greatest.

Figure 10C shows a region of the fusionzone near the HAZ. Large aggregates of awhite blocky phase (presumably ferrite)can be observed in between finer micro-constituents. Due to their unique mor-phology, these fine microconstituentswere also identified by SEM as martensite.Figure 10D depicts a narrow view of thefusion zone center. Tiny cracks (approxi-mately 50 µm long) were regularly seen,especially in the DP20 steel welds. The ir-regular topography seen in these cracksand revealed by Fig. 10E suggests shrink-age cracks, or solidification cracks; thelater being more likely, as indicated byEDS measurements of sulfur, calcium,potassium, and silicon (i.e., many ele-ments that are found in slags for deoxidiz-ing and desulfurizing steel and that arenormally removed to minimize solidifica-tion cracking tendency) (Ref. 2).

Although all our observations indicatethat the role played by microscopic crackson weld fracture is unlikely significantcompared to that of the sheet thickness,this section has brought another explana-tion for the greater susceptibility of DP20for weld interfacial fracture. These micro-scopic cracks, most distinctively encoun-tered in the DP20 steel welds, were given

further attention as chemical compositionnearby internal defects was further investigated.

Variations in Chemical Composition

For DP20, Fig. 11 shows six X-ray com-positional maps captured at a weld center,and three horizontal line scans across themaps for the major alloying elements. Inaddition to voids and cracks along theweld centerline, Fig. 11 shows that themost abundant element, manganese, wasmost distinctively microsegregated. Its re-distribution within the microstructure re-vealed a fine columnar substructure thatgrew from the fusion line to the weld cen-terline. Close examination of the map andline scan of manganese indicates periodicdepletion and accumulation every 10 to 20mm. This microstructural feature alonecan be well explained by the classicScheil’s model (Ref. 15), where the man-ganese-lean regions would be first to so-lidify and the manganese-rich regions last.Manganese profiles could thus be used totrack the solidification substructure.

The line scan of Fig. 11 shows that thepeaks for manganese, silicon, andchromium were superimposed. Unlikemanganese, chromium distribution wasquite uniform, as explained by the iron-chromium partitioning coefficient (nearlyone) (Ref. 30). Although zinc was alsofairly well dispersed throughout the fusionzone, Fig. 11 raises the possibility of zincbuildup near voids; an observation that iswell justified by zinc and iron’s limited mu-tual solubility, and zinc’s property to re-main liquid long after iron has solidified.In Fig. 11, note that the regions of high Xrays for carbon, silicon, and oxygen notonly correlated with one another, but werealso aligned with the solidification sub-structure (revealed by the microsegre-gated manganese). Another importantfeature is that all three elements wereclearly identified in the larger voids,thereby suggesting that they were proba-bly in the form of carbides and silica par-ticles left by the sample preparation. Bytracking these last three elements, micro-scopic voids could be also revealed. In Fig.10, recall that sulfur, calcium, and potas-sium were detected in microscopic voidsand cracks; an observation that first sug-gested that these elements could be linkedto the formation of internal defects. In thissection, we observe that only the contri-bution of zinc to the formation of internaldefects could not be ruled out yet.

Spot Weld Interfacial Fracture

Internal Defects

The numerous weld fracture surfaces


WELDING RESEARCH


made available by this study were finallyexamined, as a last attempt to link fracturecharacteristics, microstructure, and com-position. Figure 12A is a low-magnifica-tion secondary electron image of a typicalweld interfacial fracture in DP20. In weldswith sizeable voids, like in Fig. 12A, thepossibility that defects such as voids influ-enced fracture was not dismissed, knowingthat voids notoriously decrease spot weldproperties (Refs. 15, 31). Figure 12B is ahigh-magnification view of the largeround void of Fig. 12A. The fine dendritesat its surface present additional strong ev-idence that this void resulted from solidi-fication shrinkage. Also, the cracks at itsperimeter indicate that shrinkage causedsufficient tension to split the fine inter-locked dendrites apart; i.e., induce acracking that this time could not be linkedto a local change in chemical composition,as was found in Fig. 11. For the interme-diate region between the void and theedge of the weld, Fig. 12C shows that asubstantial part of the weld exhibits thecharacteristic microvoids of a ductile frac-ture, in spite of the high hardness values(Fig. 9) associated with the martensite mi-crostructures (Fig. 10E). Figure 12Dshows that brittle fracture by cleavage alsooccurred, particularly on the side of theweld (i.e., HAZ) where the microstructurewas the hardest — Fig. 9.

As for the zinc expulsion, the weld but-ton diameter, and the type of weld frac-ture, the effects of current, force, and timeon the shrinkage voids were all investi-gated for both steels. Quantitative resultsto compare voids in DP18 and DP20 areprovided first in Fig. 13, which like Fig. 14,was constructed after estimating pro-jected areas of all sizeable voids seen onthe interfacial fracture surfaces. Figure 13clearly shows that the percentage of thefused area covered by the voids wasgreater in DP20 than in DP18, a featurethat is also in line with the greater suscep-tibility of DP20 welds toward interfacialfracture.

Figure 13 shows that increasing theweld time, thus the weld diameter, re-duces the percentage area made by voidsin the fusion zone, and thus any contribu-tion voids might have to fracture. Figure14 confirms these results by showing thatan increase in weld diameter, as achievedby raising the current, also decreased thevoids rapidly and eliminated weld interfa-cial fracture (beyond 12 kA). Of the twoprocess variables in Fig. 14, weld force af-fected most distinctively the relative im-portance of voids. Voids were reduceddramatically at 1200 lb (5.3 kN) and 1700lb (7.6 kN), where they were nearlynonexistent. This contrasted with the 20%of the fusion zone they covered at 600 lb(2.7 kN).

Despite evidence that voids could becontrolled by process parameters, and inparticular reduced by high forces, wefound no indication that voids directlycaused interfacial fracture. However, weproved that voids and sheet thickness wererelated by showing that more voids werepresent in the thicker steel — Fig. 13.

Crack Initiation

Among all fracture surfaces, we dis-covered that those associated to abnor-mally fast weld schedules (e.g., 24 kA and5-cycle weld time) provided a new expla-nation for spot weld fracture in DP600steels. Unlike the welds examined previ-ously, the chisel testing of welds madewith high currents and extremely shortcurrent pulses produced buttons with un-common characteristics. Figure 15A–Fdetail one of those welds. First, note thepresence of several large eccentric voids,including two on the right of Fig. 15A thatappear to be superimposed. Figure 15B,for a part of the fusion zone periphery,shows an abrupt separation between aductile region with microvoids, and a den-dritic solidification structure. Typicalcharacteristics of ductile fracture are alsovisible in Fig. 15C, as are those of a cleav-age brittle fracture. Also, carbon-rich in-clusions (likely carbides) were frequentlydetected in the ductile regions, wherethey were found deep inside microvoids.Their concentrations were not quantified,since they would have improbably ex-plained DP20’s greater weld interfacialfracture susceptibility.

The microstructure shown in Fig. 15Drevealed another characteristic feature ofthese welds made with abnormally fastschedules. Figure 15D is a high-magnifi-cation view of the crack that was seen inFig. 15B. In Fig. 15D, a heavy coating ofzinc, confirmed by EDS, is seen on den-drites. Zinc was also encountered for themicroconstituent of glassy appearance ofFig. 15E. EDS measurements revealedthat its composition matched that of aniron-zinc spinel; i.e., FeZn2O4, a phase wehowever did not try to validate. In Fig. 15F,where part of the circular void of Fig. 15Ais depicted, zinc was also found quite ho-mogeneously distributed over its surface.The fact that zinc was found at dendritesand voids simply confirms that solidifica-tion cracking occurred. This cracking, bestrevealed in Fig. 15D, can be well under-stood from the binary phase diagram withiron (Ref. 30). The Fe-Zn phase diagramindicates that zinc rejection from the ironsolid solution stabilizes a zinc-rich liquidat temperatures as low as the zinc meltingtemperature; i.e., a perfect condition forcracking to occur, even under normalshrinkage conditions. Although this crack-

ing is clearly detrimental to the welds, wefound that chisel-testing zinc-infiltratedwelds had consistently produced buttons;an example that demonstrated that chiseltesting, and other mechanical tests alonewould be inappropriate for these particu-lar welds.

Conclusions

1) The two DP600 steels were success-fully resistance spot welded to producewelds resisting interfacial fracture duringchisel testing. Low currents (<10 kA), ex-tended weld times (>25 cycles), and highweld forces (>900 lb, or 4.0 kN) promotedweld button formation by producing largefusion zones and occasionally deep (butacceptable) electrode indentations. Thegreater susceptibility of the DP20 steel toproduce weld interfacial fracture was con-sidered to be the results of its thickergauge (2.0 mm) compared to the DP18steel (1.8 mm).

2) As explained by carbon equivalents,weld fusion zones in DP600 steels containedmainly martensite, and weld microstruc-tures in the DP18 steel were harder than inthe DP20 steel. The dominant effect ofsheet thickness was demonstrated by thefact that 1.8-mm-thick DP180 steel showedlower susceptibility to weld interfacial frac-ture than the thicker DP20 steel, despite itshigher carbon equivalent number.

3) Galvanized coatings generally haveinsignificant effects on weld fracture. How-ever, decreasing weld time and increasingweld current (not typical in production) re-sulted in zinc ingestion into the fusion zone,which caused solidification cracks. In thisinvestigation, ingestion of zinc was observedat the same time type of fracture changedfrom one type to the other.

4) Shrinkage voids, recognizable by theirdendritic surface morphology, were ob-served in many welds. High current, longweld time, and high weld force all helped re-duce shrinkage voids. Shrinkage voids wereless pronounced in the thinner DP18 steel,which also had lower susceptibility for weldinterfacial fracture.

Acknowledgments

The authors would like to acknowledgeChris Chen, Alexander Turley, and DavidSigler from General Motors for providingfrequent valuable discussions. Technicalsupport from Bob Cubic for the scanningelectron microscopy and Rich Valdo for themicroprobe analysis are also deeply appre-ciated.

References

1. ULSAB-AVC Consortium, TechnicalTransfer Dispatch #6 (Body Structure Materi-als), May 26, 2001.


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NOVEMBER 2005-s182 -s

2. Krauss, G. 1990. Steels: Heat Treatmentand Processing Principles, Materials Park, Ohio,ASM International.

3. Svenson, L. E., and Larsson, J. K. 2002.Steel World, 7: 21–26.

4. Marya, M., and Gayden, X. Q. 2005. Weld-ing Journal (in press) Part II of this study.

5. Rathbun, R. W., Matlock, D. K., andSpeer, J. G. 2003. Welding Journal 79(8): 207- sto 218-s.

6. Militsky, M., Pakalnins, Jiang, C., andThompson, A. K. 2003. SAE Technical Paper2003-01-0520.

7. Ghosh, P. K., Gupta, P. C., Ramavtar, A.,and Jha, B. K. 1991. Welding Journal 71(1): 7-sto 14-s.

8. Ghosh, P. K., Gupta, P. C., Avtar, R., andJha, B. K. 1990. ISIJ International 30(3):233–240.

9. General Motors Engineering Standards,Automotive Resistance Spot Welding, GM4488M, Aug. 1995.

10. Japanese Industrial Standard, JIS Z3140 — Method of Inspection for Spot Welds,2000.

11. Auhl, J. R., and Patrick, E. P. 1994. SAETechnical paper 940160.

12. Saito, T. 1983. Welding Technique 31(4): 27.

13. Takechi, H., and Akisue, O. 1985. Pro-ceedings of International Conference on HSLASteels: Metallurgy and Applications, Beijing,People’s Republic of China, pp. 977–984, ASMInternational.

14. Saito, T., and Ichiyama, Y. 1996. WeldingInternational 10(2): 117–123.

15. Olson, D. L., Edwards, G. R., Liu, S.,and Frost, R. 1993. ASM Handbook, Vol. 6,Welding, Brazing and Soldering,, MaterialsPark, Ohio, ASM International.

16. Devilliers, L., Kaplan, D., and Saint-Martin, P. S. 1986. Soudage et Techniques Con-nexes 40(11-12): 387–395.

17. Babu, S., Goodwin, G. M., Rohde, R. J.,and Sielen. 1998. Welding Journal 77(6): 249-s to253-s.

18. Lee, A., and Nagel, G. L. 1988. SAETechnical Paper 880277.

19. Quanfeng, S., Zhang, W., and Bay, N.2005. Welding Journal 84(5): 73- s to 76-s.

20. Harlin, N., Jones, T. B., and Parker, J. D.2003. Journal of Materials Processing Technology143/144: 448–453.

21. Zhang, S. 1997. International Journal ofFracture 88: 167–185.

22. Hess, W. F., and Childs, W. J. 1947. Weld-ing Journal 26(4): 712-s to 723-s.

23. Cunningham, A., Legeman, M. L., andShort, B. E. 1966. Welding Journal 45(7): 305-sto 313-s.

24. Gould, J. E. 1987. Welding Journal 67(1):1s-10s.

25. Feng, Z., Babu, S. S., Santella, M. L.,Riemer, B.W., and Gould, J. E. 1998. Proceed-ings of 5th International Conference on Trends inWelding Research, Pine Mountain, Ga., ASM In-ternational.

26. Pan, N., and Sheppard, S. D. 2003. Engi-neering Fracture Mechanics 70: 671–684.

27. Marya, M., Hector, L. G., Gayden, X.Q., and Wang, K. 2005. Paper accepted for pub-lication in Journal of Manufacturing Engineer-ing, Transactions of ASME.

28. Zhou, M., Zhang, H., and Hu, S. J. 2003.Welding Journal 62(4): 72-s to 77-s.

29. Pook, L. P. 1975. International Journal ofFracture 11: 173–176.

30. Baker, H., and Okamoto, H. 1992. ASMHandbook, Vol. 3, Alloy Phase Diagrams, ASMInternational.

31. Gear, A., Westgate, S. A., Kueze, J. C.,and Ehrstron, J. C. 1999. Welding Journal 78(3):80-s to 86-s.

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