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Introduction Maraging steels are a special class ofultrahigh-strength steels that differfrom ordinary steels because they canbe hardened by a carbon-free metallur-gical reaction (Ref. 1). The maraging18Ni steel family was first developed inthe 1960s by the International NickelCo., containing the following main com-ponents: Ni (18.9%), Co (7%), Mo(3.8%), and Fe (balance) (Ref. 2). Thisalloy presented a body-centered-cubicmartensitic structure following the qua-sibinary Fe-Ni system that is very duc-tile in the annealed condition. Itsmartensite has a high dislocation den-sity with a good response toprecipitation hardening (Ref. 3). Uponreheating at about 480ºC for 3 h, this

steel undergoes age hardening and pro-duces high yield strength, although hav-ing just a small decrease in toughness.Another important characteristic of thissteel is that martensite is formed fromthe austenite even at very low coolingrates, making it possible to treat largeparts (Ref. 2). Maraging steels have a wide rangeof applications, but they areparticularly suitable for nuclear andaerospace areas due to their excellentcombination of high yield strengthand toughness (Ref. 4). These steels donot experience significant dimensionalchanges after aging, and their low car-bon content also provides excellent di-mensional stability during the austen-ite-martensite transformation. Many studies (Refs. 5–10) had

demonstrated that the maraging steelmartensite could revert to austeniteduring aging. If enough temperatureand time are given during aging, thealloy tends to follow thethermodynamic equilibrium and a per-centage of martensite may betransformed into ferrite and austeniteby a diffusion-controlled process. Theaustenite formed this way is enrichedin nickel, since solubility of nickel ishigher in austenite than in ferrite.After cooling, the reverted austeniteremains stable if its nickel content issufficiently high. In most 18Ni maraging steels, theamount of reverted austenite isnormally not significant after usualheat treatments. Notwithstanding,austenite reversion is usually seen inthe fusion zone of welds because ofmicrosegregation in the interdendriticspaces. The elements cobalt and nickelare referred to have a low partition co-efficient in maraging steels. However,molybdenum and titanium have highsolid-liquid partition coefficients andare prone to segregate (Ref. 11).Molybdenum is known by its effect onacceleration of reverted austenite for-mation during aging (Ref. 6). The reverted austenite formed inthe interdendritic areas could not behardened by aging, having inferiorstrength and hardness compared tothe aged martensite. According to thework of Venkateswara (Ref. 12), withan 18Ni(250) maraging steel weldedby gas tungsten arc welding (GTAW),homogenization treatment is capableof eliminating the segregation effect inthe fusion zone.

Heat Treatment and Yb­Fiber Laser Weldingof a Maraging Steel

The influence of homogenization and solutionizing on an 18Ni(300) maraging steelwelded with a Yb­fiber laser source are evaluated

BY L. FANTON, A. J. ABDALLA, AND M. S. FERNANDES de LIMA

L. FANTON (leofanton@gmail.com) is with the Universidade Estadual Paulista “Julio de Mesquita Filho” (UNESP), Brazil. A. J. ABDALLAand M. S. FERNANDES de LIMA are with the Instituto de Estudos Avançados (DCTA­IEAv), Brazil.

ABSTRACT Maraging steels are iron-nickel alloys having an unusual combination of highmechanical strength and high toughness. In this work, the effects of laser welding andpostweld heat treatments have been analyzed. The fusion zone showed a cellular-den-dritic morphology with a relatively low hardness (≈450 HV) compared to the heat-af-fected zones (HAZs) and base material (≈500 HV). The segregation in the fusion zoneincreases the tendency of austenite phase formation during aging. The solutionizedand aged coupons presented the highest yield strength values (≈1890 MPa). Homoge-nization treatment showed to be effective on eliminating the as-cast dendriticstructure of the fusion zone, but decreased the yield strength to about 1350 MPa,showing that this treatment should be used carefully. The welded specimens presentedtensile strength close to the unwelded specimens, demonstrating that the ytterbium-fiber laser welding of 18Ni-type maraging steels could be considered.

KEYWORDS • Laser Beam Welding • Maraging Steels • High-Strength Steels • Precipitation Hardening

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Maraging steel has been usuallywelded with GTA and plasma arc weld-ing (PAW) with filler metal additions(Refs. 13–15). Meshram (Ref. 16)recently showed that friction stir weld-ing (FSW) also has potential for join-ing maraging steels for aerospace ap-plications. In fact, maraging steelshave demonstrated good weldability ina variety of welding techniques,including laser welding. Van Rooyen(Ref. 14) demonstrated that a Grade250 maraging steel can be welded withCO2 and Nd:YAG lasers; however, heshowed this to be very sensitive to thejoint alignment and root opening. Thehigh beam quality of the fiber lasersallows its absorption for all metal andalloys, making then a good option fordeep penetration welding (Ref. 17). The present contribution intends toevaluate the influence of the mostcommon heat treatments used inmaraging steel to an 18Ni300 alloyfabricated in the Centre of AeronauticsTechnology (Brazil). Differently from

the other steels used in the Braziliansatellite launch vehicle program,maraging steels are soft after welding,increasing the reliability of the project.These steels are being evaluated to re-place 300M and 4340 steels in parts ofthe Brazilian satellite launch vehicle.Additionally, the use of an Yb-fiberlaser source to weld maraging steelswas not reported in the literature.

Experimental Procedures The current alloy was developed bythe Centre of Aeronautics Technology inSão José dos Campos, Brazil. Sheetswith thicknesses of 2.5 and 10 mm wereused in this work, and theircompositions are shown in Table 1.Both sheets have the same origin andshould present similar composition.The differences in the compositions areprobably due to chemical analyses reso-lution and/or manufacture. No furtherinformation about the fabricationprocess was given by the manufacturer.

A 2-kW continuous-wave fiberlaser produced by IPG Co. was usedhere. The laser radiation is generatedin a 50-mm-diameter fiber doped withytterbium. The doped fiber isconnected to a process fiber 100 mmdiameter and 10 m long, which isthen connected to an Optoskand pro-cessing head. The focal length was157 mm with a minimum spot diame-ter of 100 mm. Pure argon gas at 30L/min flow rate has been used to pro-tect the sample against oxidation.The protection gas was deliveredthrough a rounded copper tube of 3mm internal diameter directly overthe irradiated area. The welding headstaying still and the samples movingare realized by a CNC table. Figure 1shows the laser experimental weldingsystem used. Initial welding tests have beencarried out on 10-mm-thick platesusing different combinations of beampower and welding speed. The sampleswere previously ground to guarantee aflat and homogeneous surface on eachsample. Just before welding, the sam-ples’ surfaces were washed with

Fig. 1 — Photo of the laser welding system used.

Table 1 — Chemical Composition of the 18NiMaraging Steel

Element Content (%) Content (%)(Alloy A) (Alloy B)

C 0.008 0.01Ni 19.12 17.85

Mo 4.94 4.96Co 9.66 9.32Al 0.089 0.14Ti 0.77 0.8Cu 0.076 —Ca — 0.04Zr — <0.01S 0.002 0.007P 0.008 0.002Si 0.07 0.05

Mn 0.011 0.01Cr 0.043 0.04

Table 2 — Welding Parameters for Initial Tests

Sample Laser Scanning Speed Beam Power(cm/min) (W)

1 300 14002 300 16003 120 16004 120 14005 120 18006 180 18007 240 1800

8 (pre­aged) 120 18009 (pre­aged) 180 180010 (pre­aged) 240 180011 (pre­aged) 300 1800

Fig. 2 — Tensile test specimen dimensions.

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ethanol. Welding parameters for theseinitial tests are shown in Table 2. Theobjective of these tests is to give somedirections on choosing good weldingparameters for the fabrication oftensile specimens. To avoid joint align-ment problems, the bead-on-plate tech-nique was chosen for the fabrication oftensile specimens. The effects of aging,homogenization, and solutionizingwere investigated. The aging treatmentwas performed at 480°C during 3 h. Forhomogenization, samples were heatedat 1150°C for 1 h, and the solutionizingtreatment was done at 815°C for 1 h.The samples were air cooled to roomtemperature after each heat treatment. The microstructure of the weldedsamples were analyzed by optical mi-croscopy (Zeiss Epiphot 2000) usingthe following etchants: 1) modifiedFry’s reagent (150 mL H2O, 50 mLHCl, 25 mL HNO3, and 1 g CuCl2), 2)

Nital 15%, and 3)sodium metabisul-fite solution (10 gin 100 mL H2O). Microhardness measurements werecarried out using a Vickers microhard-ness tester (Future-Tech FM-700) witha 300 gf load for 10 s. Measurementshave been done in a traverse cut of theweld bead at 0.5 mm from the weldface and 0.1 mm betweenindentations. For the fabrication of the tensilespecimens, bead-on-plate welds wereperformed on plates of approximately2.5 mm thickness. Based on previousexperiments, it used a laser scanningspeed of 180 cm/min and laser powerof 1800 W. The laser beam wasfocused on the sample surface. Speci-mens for tensile tests were cut outfrom the welded plates, according toFig. 2. The weld bead was oriented

transversely to specimen’s length. The tensile specimens were dividedin four groups with different postweldheat treatment conditions as follows: Group A: homogenized,solutionized, and aged Group B: homogenized and aged Group C: solutionized and aged Group D: aged. Homogenization treatment createdan oxidized layer. These samples wereground using SiC 600 paper until theoxidation was visibly removed. At leastthree specimens of each heattreatment condition were tested andunwelded specimens of each groupwere also tested for comparison. Theexistence of pores in the weld beadwas verified by X-ray analysis, andsamples with visible pores wererejected. Tensile tests were carried out

Fig. 3 — Different welding parameters. A — Weld bead depth; B — width.

Fig. 4 — Cellular/dendritic solidification structure of the fusionzone for a nonaged sample (Alloy A). Etched with Nital 15%.

Fig. 5 — Optical microscope image showing the fusion zone andHAZ for a homogenized sample (Alloy B). The fusion zone nolonger shows the as­cast dendritic structure. Etching: modifiedFry’s reagent.

A B

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on an Instron 338 machine with aloading rate of 1 mm/min. Thefracture surfaces were observed underscanning electron microscope (SEM).

Results and Discussion

Initial Welding Tests

The measured weld beads depths andwidths are shown in Fig. 3 for the 10-mm-thick plates. The samples that wereaged before welding (pre-aged) showedsimilar dimensions to the unaged speci-mens using the same parameters, indi-cating that the precipitates have a lowinfluence on the laser absorption. Thetensile specimens used in this workwere not aged before welding.

The initial results provided an idea

about the behavior ofthe alloy after Yb-fiber laser welding.Based on theseresults, some testshave also been donewith the 2.5-mm-thick sheet, consider-ing that the heat flowis different in a thin-ner sample. The pres-ence of pores andweld depth were ana-lyzed. The bestresults were obtainedusing a weld speed of180 cm/min andbeam power of 1800W. These were thenthe parameters usedfor the tensile speci-mens.

Microstructure and Hardness

The fusion zone exhibits a cellular/dendritic morphology, as shown in Fig.4. After the homogenization treatment,the fusion zone no longer presented thecellular/dendritic structure, as shownin Fig. 5. This result indicates that thehomogenization treatment is effective,as expected, on eliminating the segre-gation of the fusion zone. Figure 6 shows the fusion zone ofan aged sample as well as asolutionized and aged sample etchedwith sodium metabisulfite solution.The white regions revealed a similarpattern found by Venkateswara (Ref.12) using an electron probe microana-lyzer (EPMA) about the distribution of

molybdenum, which is reported tosegregate interdentrically. Themetabisulfite solution showed to be agood option to reveal themolybdenum-rich areas. Near to the fusion zone in the as-welded condition, there is an austeni-tized heat-affected zone (HAZ 1)followed by an aged zone (HAZ 2), asshown in Fig. 7. The metallurgical be-havior of the HAZ can be understoodthrough the analysis of hardnessacross the weld joint. Figure 8A shows the hardness pro-file for the sample welded with 1800 Wlaser power and a welding speed of 180cm/min. In the as-welded condition,the fusion zone and HAZ 1 have simi-lar hardnesses of about 300 HV. Thehardness values increase in the HAZ 2and then decrease as long as it gets far-ther from the weld interface. The age-hardening effect observed in HAZ 2 isproportional to the heat input, as ob-served in Fig. 8B, where a relatively lowwelding speed of 30 cm/min was used.In this case, the hardness valuesobserved on HAZ 2 were relativelyhigh. All other welding parametersused in this work presented the samegeneral characteristics. The temperatures experienced byHAZ 1 were high enough to transformit into austenite. Although this trans-formation occurs at temperaturesaround 700ºC, the austenite will onlybe transformed back to martensiteduring cooling at much lower tempera-tures, where no aging effect isobserved, preserving the original softmartensite and hardness values simi-lar to the base material.

Fig. 6 — Fusion zone images. A — The aged condition; B — solutionized and aged condition. Alloy B, Etching: sodium metabisulfite solution.

Fig. 7 — Fusion zone image and the two heat­affected zones(HAZ) (Alloy A). Etched with Nital 15%. HAZ 1 was austenitizedand HAZ 2 was aged during welding.

A B

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The HAZ 2 had not beentransformed into austenite, but thetemperatures achieved were highenough to aging, even in such a smallperiod of time, increasing itshardnesses values. The higher thetemperature achieved, the higher theaging effect. The shadow observed inHAZ 2 is actually caused by the precip-itates generated during aging. After the postweld aging treatment,the two HAZs achieve similar hardnessof about 500 HV (Fig. 8A, B). In thiscondition, the fusion zone presentedlower hardness values, probablycaused by the effects of segregationand formation of reverted austenite. The homogenization treatment afterwelding was not able to make the fusionzone achieve the same hardness valuesof HAZ and base material. Figure 9shows the hardness profile of a homog-enized, solutionized, and aged sample.

Tensile Strength Tests

Figure 10 shows the results of ulti-mate tensile strength, yield strength,and total elongation for all heat-treated conditions. Group C(solutionized and aged) had thehigher strength values, followed bygroup A (homogenized, solutionized,and aged), group D (aged), and groupB (homogenized and aged). Figure 11 shows a traverse cut of thefractured region with arrows indicatingthe fractured surfaces. The specimensfractured preferentially at the fusionzone. In some cases, the crack wentthrough the fusion zone and also theHAZ, as exemplified by Fig. 11B. The welded specimens of all groupspresented quite close mechanical prop-

erties whencompared to the un-welded specimens,indicating that thematerial has goodweldability. The lowhardness of the fu-sion zone did nothave a noticeable in-fluence on the ten-sile strength. Thewelded specimenspresented inferiorelongationcompared to the un-welded specimens,except for group B,where the averagevalue was higher.Nevertheless, it isimportant toconsider that the in-ferior hardness of the fusion zonecould cause deformation to berestricted to this small area. Homogenization treatment wasdetrimental to mechanical propertiesin all cases. One of the reasons,pointed out by literature, is the grainsize growth after homogenizationtreatment (Refs. 12, 18, 19) since theHall-Petch relationship applies (Ref.20). Moreover, Rack (Ref. 20) showedthat 18Ni maraging steels with a largegrain size may fail catastrophically dueto the propagation of intergranularcracks. According to Venkateswara(Ref. 12), solutionizing heat treatment(in the same temperature and timeused in this work) was able to reducegrain size after homogenization treat-ment and restore some mechanicalproperties. For the alloy used in thiswork, solutionizing treatment after

homogenization (Group A) was able torecover some strength, compared tospecimens that were onlyhomogenized (Group B), but at stilllower values when compared to speci-mens that were only solutionized(Group C). The detrimental effects of homoge-nization observed in this work, evenafter solutionizing treatment, mighthave been intensified by the oxidationlayer present on the specimen’ssurface. Even after grinding, anoxidized layer of about 100 mm couldstill be observed by opticalmicroscope. Figure 12 shows atraverse cut image of the fracturedsurface of a homogenized specimen,where the oxidized layer ishighlighted. It is possible to see someoxidation points where the fracturemay have been originated.

Fig. 8 — Hardness distribution across weldment of an as­welded (white dots) and aged sample (black dots). Welding parameters are asfollows: A — 1800 W and 180 cm/min; B — 1800 W and 30 cm/min.

Fig. 9 — Hardness distribution across the weldment of a homog­enized, solubilized, and aged sample welded with 1800 W laserpower and 60 cm/min scanning speed. The fusion zone still pres­ents lower values of hardness compared to the HAZ and basematerial even after homogenization treatment.

A B

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An intergranular crack at theoxidized region can be observed on thefracture surface image of Fig. 13.

Conclusions The stress strength values of thewelded specimens were close to thatobtained for sheet specimens, demon-strating that the alloy studied in thiswork was successfully welded by theytterbium-fiber laser source. It was observed there were two dif-ferent HAZs — an austenitized zoneand an aged zone. The transformationthar occurred in the HAZ did not havea significant influence on the tensilestrength since the fracture occurredpreferentially in the fusion zone. Among all heat treatmentconditions used in this work, the solu-tionized and aged specimens showed

the best yield strength values, around1890 and 1900 MPa for welded andsheet specimens, respectively. The ho-mogenized samples, on the otherhand, had their mechanical propertiesseverely affected. The oxidized layerformed on the homogenizedspecimens appears to be the origin ofcrack formation. The fusion zone exhibits acellular/dendritic morphology wherethere is segregation of alloy elements,especially titanium and molybdenum.The sodium metabisulfite solutionetching reagent was capable of reveal-ing the molybdenum-enriched areas ofthe as-cast metal. This segregation in-

creases the tendency of revertedaustenite formation on the fusionzone, but the strength values obtainedin this work were not severely affectedby this phenomenon. Despite theexpected effect of homogenizationtreatment on recovering the hardnessvalues of the fusion zone, the alloy usedin this work presented lower hardnessvalues in the fusion zone compared tothe HAZ and base material. The oxidation effect ofhomogenization reported in this workindicates that the use of anatmosphere-controlled furnace shouldbe necessary. Another problem alreadyreported in the literature about

Fig. 10 — Tensile test results. A — Ulti­mate tensile strength (UTS); B — yieldstrength (YS); C — elongation.

Fig. 11 — Traverse cut images of the fractured surface for specimens of groups A to D.The arrows indicate the fractured region. Etching: modified Fry’s reagent.

Fig. 12 — Traverse cut image of the fractured surface of a homogenized specimen. Thearrow indicates where the fracture began. It is possible to observe some points wherethe oxidation was more intense. Etching: modified Fry’s reagent.

AA B

B C

C

D

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homogenization treatment of marag-ing steels is the intergranular crackcaused by grain growth. Consideringthat the tensile strengths forunwelded and welded specimens arealso very close, this work suggests thatthe use of homogenization treatmentas a way to eliminate the segregationeffects in the fusion zone may not beworthwhile and should be evaluatedseparately for each maraging alloy andwelding technique.

The authors thank CAPES, Pro-De-fense Project 014/08, the Instituto deEstudos Avançados (IEAv/DCTA), andDr. Rogério de O. Hein from UNESP.

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Fig. 13 — Images of the fracture surface of a homogenized and aged specimen. A — Transition between the fragile oxidized layer and the basemetal; B — a highlighted region where an intergranular crack occurred. The oxidized region shows a fragile rupture while the base metalshows the expected ductile fracture mode with a presence of dimples.

Acknowledgments

References

A B

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