Astm stp 1532 2012

Jin Zhang,1,2and Dong-Hua Yang3Effect of Aluminum Coating by Magnetron Sputtering on Corrosion Resistance of AZ31B Alloy ABSTRACT: In this study, aluminum coatings were prepared by dc ma

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A Stojko, M F Hansen, J Slycke, and M A J Somers 44 Microstructural Characterization of an AISI D2 Tool Steel Submitted to Cryogenic

Treatment

P F da Silva Farina, C A Barbosa, and H Goldenstein 57

Diffusion Processes

Interdiffusion Phenomena of Zirconia-Nitride Layers on Coated AISI 310 Steel

B J Gómez, L Nachez, R Caruso, A Díaz-Parralejo, J N Feugeas,

The Effect of Si on the Precipitates for Nb Bearing High Purity Ferritic Stainless Steel

M Ito, Y Honda, M Aramaki, Y Kato, and O Furukimi 107 Effect of the Thermo-Mechanical Processing Characteristics on the Recrystallization

Steel AISI 444 Using Austenitic Stainless Steels Filler Metals

P D Antunes, E O Corrêa, N D Barbedo, P de Oliveira Souza, J L Gonçalves,

and A A Diacenco 176 Modifi cation of NiAl Intermetallic Coatings Processed by PTA with Chromium

Carbides

D H S Yano, C Brunetti, G Pintaude, and A S C M D’Oliveira 190 Morphology of Precipitates in a First Stage Low Pressure Turbine Blade

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N Miura and Y Kondo 205

Modeling and Simulation

Mathematical Modeling and Computer Simulation of Fatigue Properties of Quenched and Tempered Steel

B Smoljan, D Iljkic´, F Traven, and J Mrša 221 Modeling of Hardness of Low Alloy Steels by Means of Neural Networks

R N Penha, L F Canale, and A C Canale 234 Simulation of Austenitization Processes in Fe–C Steels by Coupled Cellular

Automaton and Finite Difference Methods

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New Materials

New Cast Iron Alloys with High Wear Resistance at Elevated Temperatures

W Theisen and G Gevelmann 275 Heat Treatments of Fe-Mn-Si Based Alloys: Mechanical Properties and Related Shape Memory Phenomena

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M J Kupczyk, A Michalski, P Siwak, and M Rosinski 313 Effect of Gas Pressure on Active Screen Plasma Nitriding Response

A Nishimoto, K Nagatsuka, R Narita, H Nii, and K Akamatsu 327

M Bao, X Xu, R Liang, L Yu, and H Sun 383 Infl uence of Structure on Brittleness of Boron Nitride Coatings Deposited on 

Cemented Fine-Grained Carbides

M J Kupczyk and P Siwak 392 Impact Wear Performance of Thin Hard Coatings on TiC Cermets

R Veinthal, F Sergejev, C E Yaldiz, and V Mikli 404

Quenching and Quenchants

Hot-Dip Galvanizing Process Using ZinQuench for Processing Advanced

High-Strength Steels

F Huber and W Bleck . 423 Dependence of the Heat Transfer Coefficient at Quenching on Diameter of Cylindrical Workpieces

B Lišcˇic´, S Singer, and H Beitz 438 Effect of Antioxidants on Oxidative Stability and Quenching Performance of Soybean Oil and Palm Oil Quenchants

G Belinato, L C F Canale, and G E Totten 450 Infl uencing Factors of Heat Transfer Coefficient in Air and Gas Quenching

B Xiao, G Wang, R D Sisson, Jr., and Y Rong 472 Effect of Cooling Rate During Quenching on the Toughness of High Speed Steels

C S Gonçalves, A L Slaviero, R A Mesquita, A P Tschiptschin,

and P de Tarso Haddad 484

Surface Hardening

Surface Hardening of an AISI D6 Cold Work Steel Using a Fiber Laser

F A Goia and M S F de Lima 499 Effect of Temperature and Pressure on Wear Properties of Ion Nitrided AISI 316 and

409 Stainless Steels

F A P Fernandes, S C Heck, R G Pereira, P A de Paula Nascente,

and L C Casteletti 512 Infl uence of Surface Hardening Depth on the Cavitation Erosion Resistance of a Low Alloy Steel

S Goulart-Santos, R D Mancosu, C Godoy, A Matthews, and A Leyland 524

Cracking Resistance and Impact Wear of Thin and Thick Hard Coatings Under Cyclic Loading

P Kulu, M Saarna, F Sergejev, A Surženkov, and A Sivitski 574 Author Index 591 Subject Index 595

This special issue assemblies 40 papers presented in this conference showing important fi ndings regarding: New Materials, Quenchants and Quenching, Tribology, Modeling and Simulation, Plasma Technology, PVD and CVD, Materials Characterization, Cryogenic Treatments, Surface Hardening, Corrosion, Diffusion Process and Equipments

Our special thanks to all the authors and co-authors for their excellent contribution promoting a successful conference permitting also the edition

of this STP

CORROSION

Jin Zhang,1,2and Dong-Hua Yang3

Effect of Aluminum Coating by Magnetron

Sputtering on Corrosion Resistance

of AZ31B Alloy

ABSTRACT: In this study, aluminum coatings were prepared by dc magnetronsputtering on AZ31B magnesium alloy The influence of sputtering parameters(include sputtering current, argon pressure, and deposition time) on corrosionbehavior was investigated by potentiodynamic polarization tests in 3.5 % NaClsolution The corrosion morphology was examined in detail by scanning electronmicroscopy and optical microscopy, respectively It was found that the corrosioncurrent density of magnesium alloy with aluminum coating was 2–3 orders ofmagnitude less than that of bare alloy The corrosion potential with aluminumcoating had been positive shift The corrosion resistance of the coatings wasstrongly affected by its structure and residual stress, which depended on the pro-cess condition Severe corrosion will occur after the aluminum coating is dam-aged due to the interaction of galvanic corrosion and other corrosion forms

KEYWORDS: AZ31B magnesium alloy, magnetron sputtering, aluminumcoating, corrosion resistance

Introduction

As the lightest of structural metals, magnesium alloy has an excellent processingperformance, good biocompatibility, high thermal conductivity, high dimensionalstability, high damping, and good electromagnetic shielding characteristics These

Manuscript received July 8, 2010; accepted for publication March 3, 2011; published online April 2011.

Copyright V C 2011 by ASTM International, 100 Barr Harbor Drive, PO Box C700, West Conshohocken, PA 19428-2959.

3

properties make it valuable applied in many fields such as aerospace, automotive,electronics, telecommunications, and even used as an implant metal However,poor corrosion resistance limits its widespread use in many applications [1] Inaddition, magnesium cannot form a self-healing passive surface contrast to alumi-num and titanium because there is a misfit between the hydroxide lattice in thesurface region and the lattice of the bulk material Therefore, the magnesium hy-droxide film is loose and it is undermined by the corrosion process [2] In order toimprove the corrosion resistance of magnesium alloy, proper surface treatments

to produce anti-corrosion protection coatings on the substrate become essential.New technology such as plasma electrolytic oxidation, metal coating (e.g., alumi-num), physical vapor deposition (PVD), and ion implantation has become thefocal point of worldwide attention in recent years

Because aluminum and its alloy have the ability to self-repair against sion and Al element is one of the main compositions of Mg alloy, aluminumcoating is designed to protect magnesium alloy from corrosion resistance Alu-minum coating has been given more concern in the past years There are manyways to prepare Al coating, such as thermal spray [3–5], PVD [6–8], chemicalvapor deposition [9], electro-deposition [10,11], and laser cladding [12] Wu

corro-et al [6] prepared pure Al and Ni thin films successfully by rf-sputtered on

a low cohesive strength with the substrate and thus will influence the corrosionresistance Wu et al [7,13] developed Al coating, aluminum/titanium multilayercoating, and Ti–Al–N/Ti–Al duplex coating on AZ31 alloy by multi-magnetronsputtering; most of these coatings could improve the corrosion resistance ofmagnesium alloy to some extent

In the present work, aluminum coatings were prepared by magnetron tering on AZ31B magnesium alloy The influence of sputtering parameters(include sputtering current, argon pressure, and deposition time) on corrosionbehavior were investigated by potentiodynamic polarization tests

sput-Experimental

Substrate Material

Rolled AZ31B magnesium alloy substrates ð25  25  1 mmÞ were used in thisstudy The composition of the AZ31B substrates was 2.5–3.5Al, 0.20Mn,0.6–1.4Zn, 0.04Cu, 0.10Si, 0.005Ni, 0.005Fe, and balance Mg The substrates

the surface oxides and then polished up to 1000# SiC emery paper Then, theywere cleaned in acetone for 10 min and in ethanol for 10 min by ultrasoniccleaning, respectively

system) with an Al (99.99 %) target placed at a distance of 60 mm from the strate Argon ion cleaning before deposition was carried out to avoid contami-nation and improve adhesion The deposition process was performed at roomtemperature without any additional heating to the substrate The base pressure

The sputtering current, argon pressure, and deposition time of the ing system were chosen as the main process parameters and the Al coating wasdeposited by varying these parameters The deposition experiments includethree parts, the detail parameters are listed in Table 1 The first four runs wereperformed at different sputtering currents of: 0.82 A (D1), 0.92 A (D2), 1.22 A(D3), and 1.62 A (D4) while keeping argon pressure at 0.5 Pa and depositiontime at 150 min; the second five runs were performed at different argon pres-sures of: 0.3 Pa (Q1), 0.4 Pa (Q2), 0.5 Pa (Q3), 0.7 Pa (Q4), and 0.9 Pa (Q5), fixingthe sputtering current constant at 1.02 A and deposition time at 150 min; thethird three runs were performed at different deposition times of: 120 min (S1),

sputter-300 min (S2), and 330 min (S3), keeping the following parameters constant forsputtering current: 1.02 A and argon pressure 0.5 Pa

Corrosion Measurements and Coating Characterization

Corrosion studies were made in 3.5 % NaCl solution using a potentiodynamicpolarization unit (M273, EG&G) consisting of three electrodes One of the elec-trodes served as the working electrode and a platinum electrode served as thecounter electrode, which was kept parallel to the working electrode A saturatedcalomel electrode (SCE) with a capillary acted as the reference electrode Thiselectrode was kept near the surface of the working electrode and measured thepotential of the working electrode

area exposed The samples were immersed in the solution for 1=2 h so that asteady state equilibrium potential, known as the open circuit potential, was

TABLE 1—The parameters for preparation of the Al coating.

Sample Sputtering Current [A] Argon Pressure [Pa] Deposition Time [min] D1 0.82 0.5 150

attained The polarization of the samples was carried out first cathodically andthen anodically at 1 mV/s at room temperature Current and potential were

method The specimen was taken out and ultrasonically washed in pure alcoholafter corrosion test Then, its surface morphology was observed through scanningelectron microscopy (SEM, JSM-6460LV) operating at 20 kV The surface mor-phology was also examined by optical microscopy with a magnification of 10.Results and Discussion

Corrosion Behavior in 3.5 % NaCl Solution

The potentiodynamic polarization curves obtained for Al coating samples underdifferent deposition parameters in 3.5 % NaCl solution are shown in Fig 1 Thepolarization behavior of the bare AZ31B is also depicted in Fig 1, and it hadregistered a corrosion potential of 1.527 mV versus SCE, and a corrosion cur-

after coating aluminum during the anodic polarization The corrosion potential

the resulting data are summarized in Table 2 The effect of different processes

The crystallographic and geometric structure as well as internal stressdepended on the process condition Figure 2 shows that the corrosion potential

of the sample deposited with different sputtering currents (a), argon pressure (b),and deposition time (c), respectively As shown in the figure, the corrosion poten-tial of samples were positive shifted to 1.442 (D2), 1.463 (Q6), and 1.361 V(S2), respectively, after being deposited with different aluminum films In addi-tion, both anodic and cathodic current densities of the magnesium alloy after

of bare alloy The test results indicate that the corrosion resistance of Al coatedAZ31B alloy have been improved, however the degree of improvement was quitedifferent because the properties of the coating were strongly affected by the sput-tering parameters

As shown in the photograph by optical microscopy (Fig 2(d)), the corrodedsurfaces after polarization tests also revealed the differences in the corrosiondamage of the Al coatings It can be seen that the surface of the bare AZ31B sub-strate showed pitting corrosion When Mg alloy is in a corrosive medium, po-

surface The larger magnification micrograph shown in Fig 3(a) demonstratedthat the surface of the pitting area was completely covered by a thick and

Compared to the bare substrate, the surface of the coated samples did notshow severe corrosion and the corrosion rates slowed down Figure 2(d) showedthat there were many micro cracks and a pinhole (The higher magnification

6 JAI  STP 1532 ON 18TH IFHTSE CONGRESS

micrograph shown in Fig 3(b)) that appeared in the surface, owing to the

cracks and pinholes not only exacerbated the corrosion rate, but also causedstress concentration There were numerous similar defects in coatings when thesputtering current was 0.82 A With an increase of sputtering current, theamount of micro crack decreases However, a large single crack occurred aftercorrosion test as the internal stress and properties of samples did not match(shown in Fig 3(c)) When the sputtering current increased to 1.6 A, there was

no significant corrosion damage in the D4 sample except a few corrosion spots.With increasing sputtering current, argon ions from the plasma become moreenergetic They can hence release higher energies to aluminum atoms duringthe sputtering process The sputtered atoms have consequently left the target

FIG 1—Polarization curves of samples (a) different sputtering current, (b) differentargon pressure, and (c) different deposition time on AZ31B in 3.5 wt % NaCl solution

with higher kinetic energy Therefore, sufficient high sputtering current isessential for improving the corrosion resistance of Al coated AZ31B.

Proper argon pressure is also an important factor to the quality of Al ing At lower pressure, there are not enough atoms to bombard the target,whereas sputtered Al atoms experience more collisions, reach the substrate atrelatively lower energy, and hence decrease the probability of depositing.Figure 2(b) and 2(d) showed that the Al coating did not reveal a significant pro-tective effect at lower and higher argon pressure, which can be ascribed to theloose and numerous defects in coatings The potential Ecorr of Al coated alloywould not increase until the deposition argon pressure reached 0.7 Pa com-pared with bare Az31B Additionally, there is no large crack except a few corro-sion spots on the eroded surface

coat-With deposition time increase, the corrosion resistance improved cantly, especially for deposited 300 min aluminum coating (S2); the corrosionpotential increased by 166 mV than the bare magnesium alloy Meanwhile, cor-rosion current density also decreased by three orders of magnitude There wasalso no corrosion damage found on the surface, the same as that of before cor-rosion tests In our previous work [14], it was found that the thickness and criti-cal load of Al coatings varied greatly with deposition time, and leading to aneffect on the corrosion resistance of samples Generally, longer deposition timecan increase coating thickness; however, the growth of coatings will slow downwhen the deposition time reaches a critical value Furthermore, a longer deposi-tion time would increase the internal stress between the Al coating and AZ31B,

penetrated into the interface combined the effect of stress and accelerated thecorrosion rate Ultimately, the corrosion became worse

TABLE 2—Polarization data of Al coating on AZ31B.

Sample E corr [V versus SCE] I corr ½lA  cm 2  AZ31B 1.527 8:91  10þ01D1 1.487 5:44  10þ00D2 1.442 2:48  10þ00D3 1.486 4:10  10þ00D4 1.469 4:75  10þ00Q1 1.545 8:64  10 þ00

8 JAI  STP 1532 ON 18TH IFHTSE CONGRESS

FIG 3—SEM micrographs of (a) corrosion pit of AZ31B substrate, (b) a pinhole, and(c) a crack of Al coated sample after corrosion testing.

10 JAI  STP 1532 ON 18TH IFHTSE CONGRESS

Corrosion Mechanism of Al Coating on AZ31B Alloy

NaCl solution is of heavy erosion to Mg alloy; even small amounts of chlorideion usually break down the protective film on magnesium Pitting corrosion

3.5 % NaCl solution On the basis of analyzing the above results, the corrosionmechanism of Al coated AZ31B magnesium alloy in 3.5 % NaCl solution wasput forward schematically in Fig 4 First, after coated aluminum, the corrosion

reflected by the passive-like behavior This is a type of corrosion initiated at thesurface of Al coating in the presence of aggressive chloride ions according to thereactions

the interface of Al coating/Az31, thus leading to the formation of new corrosive

FIG 4—Schematic representation of the corrosion process of Al coated AZ31B uponimmersion for 3.5 % NaCl solution

products MgðOHÞ2 Moreover, dissolution of the a-Mg grains and Mg17Al12

phase within the lamellar aggregate happen on substrates according to the lowing reactions:

form at the interface, according to reaction Eq 6

on Al coatings which has intrinic stress prepared by sputtering, thus resulted inthe lifted/damaged of the coating (Fig 3(c)) In addition, where hydrogen evolu-tion happened, the film would rupture to form cracks at the force of the increas-

If the corrosion processes continue, the reduction reaction will cause localalkalization around cathodic layers/particles Aluminum oxide is not stable insuch an environment; the alkalization may happen in thin Al films [15] Severecorrosion will occur due to the interaction of galvanic corrosion and other cor-rosion forms Therefore, it is important to obtainin thicker, denser, and free-defect Al coating by optimizing the deposition parameters in order to improvingthe corrosion resistance of AZ31B magnesium alloy

Conclusions

Aluminum coatings were deposited on AZ31B magnesium alloy using dc netron sputtering The influence of sputtering parameters on corrosion behav-ior was different The coated sample demonstrated good corrosion resistance in3.5 % NaCl solutions than that of the bare substrate It was found that the corro-

alu-minum coated samples had been positive shift A sufficiently high sputteringcurrent (1.6 A) and proper argon pressure (0.7 Pa) are essential for improvingthe corrosion resistance of Al coated AZ31B With deposition time increase, thecorrosion resistance improved significantly Severe corrosion will occur afterthe aluminum coating is damaged

Acknowledgments

The writers gratefully acknowledge the Key Research Projects of the Ministry ofEducation, China (Grant No 2070950) and the Beijing Key Laboratory for Cor-rosion, Erosion and Surface Technology for their financial support to thisresearch

12 JAI  STP 1532 ON 18TH IFHTSE CONGRESS

on the Interfacial Characteristics of Aluminum Coated AZ91D Magnesium Alloy,” Mater Sci Forum, Vols 546–549, 2007, pp 529–532.

[4] Zhang, J and Sun, Z F., “Research on Magnesium Alloy AZ91D Surface Coating by Metallizing Aluminim,” China Mech Eng, Vol 13, 2002, pp 2057–2059 (in Chinese) [5] Zhang, J and Wang, Y., “Effect of Heat Treatment on Microstructures and Proper- ties of Zinc-Aluminum Coating on AZ91D Magnesium Alloy,” Key Eng Mater., Vols 373–374, 2008, pp 55–58.

[6] Wu, S K., Yen, S C., and Chou, T S., “A Study of r.f.-Sputtered Al and Ni Thin Films

on AZ91D Magnesium Alloy,” Surf Coat Technol., Vol 200, 2006, pp 2769–2774 [7] Wu, G., “Fabrication of Al and Al/Ti Coatings on Magnesium Alloy by Sputtering,” Mater Lett., Vol 61, 2007, pp 3815–3817.

[8] Bohne, Y., Manova, D., Blawert, C., Stormer, M., Dietzel, W., and Mandl, S.,

“Influence of Ion Energy on Morphology and Corrosion Properties of Mg Alloys Formed by Energetic PVD Processes,” Nucl Instrum Methods Phys Res B, Vol.

257, 2007, pp 392–396.

[9] Liu, Y., Overzet, L J., and Goeckner, M J., “Chemical Vapor Deposition of Aluminum from Methylpyrrolidine Alane Complex,” Thin Solid Films, Vol 510, 2006, pp 48–54 [10] Zhang, J., Yan, C., and Wang, F., “Electrodeposition of Al-Mn Alloy on AZ31B Mag- nesium Alloy in Molten Salts,” Appl Surf Sci., Vol 255, 2009, pp 4926–4932 [11] Chang, J K., Chen, S Y., Tsai, W T., Deng, M J., and Sun, I W., “Electrodeposition

of Aluminum on Magnesium Alloy in Aluminum Chloride midazolium Chloride (EMIC) Ionic Liquid and Its Corrosion Behavior,” Electrochem Commun., Vol 9, 2007, pp 1602–1606.

(AlCl3)–1-Ethyl-3-Methyli-[12] Chen, C., Wang, M., Wang, D., Jin, R., and Liu, Y., “Lser Cladding of Al+Ir ders on ZM5 Magnesium Base Alloy,” Rare Metals, Vol 26, 2007, pp 420–425 [13] Wu, G., Wang, X., Ding, K., Zhou, Y., and Zeng, X., “Growth and Corrosion of Alu- minum PVD-Coating on AZ31 Magnesium Alloy,” Mater Charact., Vol 62, 2009,

OPow-pp 4325–4327.

[14] Zhang, J., Yang, D., and Ou, X., “Microstructures and Properties of Aluminum Film and Its Effect on Corrosion Resistance of AZ31B Substrate,” Trans Nonfer- rous Met Soc China, Vol 18, 2008, pp s312–s317.

[15] Wu, G., Zeng, X., and Yuan, G., “Growth and Corrosion of Aluminum PVD-Coating

on AZ31 Magnesium Alloy,” Mater Lett., Vol 62, 2008, pp 4325–4327.

Milan Maroˆnek,1Jozef Ba´rta,2and Katarı´na Ba´rtova´3

Comparison of the Laser and Electron

Beam Welding of Steel Sheets Treated

by Nitro-Oxidation

ABSTRACT: Nitro-oxidized steel specimens welded by a CO2laser and anelectron beam were examined for quality and integrity by visual inspection,microstructure, and microhardness analysis Visual inspection of the speci-mens welded by laser revealed the lack of root penetration at maximal weld-ing speed from the selected welding speed range (30 to 60 mm=s) and aneat weld at the welding speed of 40 mm=s However, in comparison to laserbeam welding, the electron beam welding (EBW) process produced the weldjoints with a wider heat affected zone, while the weld joints were much moreprone to porosity The microstructure of the weld joints made by both thelaser beam welding and EBW methods did not show any abnormalities inphase composition owing to the presence of nitrides in the surface layer ofthe materials welded after nitro-oxidation

KEYWORDS: nitro-oxidation, laser beam welding, electron beam welding

Introduction

Surface treatment by nitro-oxidation is generally used for improving the sion resistance and mechanical properties of materials [1] It provides a low-cost alternative to using very expensive materials, especially in the case of bulky

corro-Manuscript received September 14, 2010; accepted for publication October 18, 2011; published online November 2011.

1 Professor, PhD., Faculty of Materials Science and Technology, Slovak Univ of Technology, Trnava 91725, Slovak Republic.

18th IFHTSE Congress, July 26–30, Rio de Janeiro, Brazil.

Cite as: Maroˆnek, M., Ba´rta, J and Ba´rtova´, K., “Comparison of the Laser and Electron Beam Welding of Steel Sheets Treated by Nitro-Oxidation,” J ASTM Intl., Vol 9, No 2 doi:10.1520/JAI103375.

Copyright V C 2011 by ASTM International, 100 Barr Harbor Drive, PO Box C700, West Conshohocken, PA 19428-2959.

14

Reprinted from JAI, Vol 9, No 2 doi:10.1520/JAI103375 Available online at www.astm.org/JAI

machine parts [2] On the contrary, the welding of surface treated materialsposes several problems to welding engineers First, the surface layer may nega-tively influence the stability of the arc and beam welding processes, having asubsequent negative impact on the weld joint quality [3,4] Another importantproblem is related to the surface layer deterioration during fusion welding.Owing to the welding process, the material is exposed to the strong thermaleffect, and the surface layer loses its primary function Regarding the previouslymentioned problems, the welding processes characterised by a narrow weldbead and heat affected zone are obviously an advantage The research and pub-lications concerning the welding of steel sheets treated by nitro-oxidation arenot abundant Our previously published papers have pointed out the difficultiesemerging in GMAW and resistance welding of nitro-oxidized steel sheets [5,6].This paper compares two welding methods using the high energy density, i.e.,the methods of laser and electron beam welding with regard to the weld jointquality and the minimization of surface layer deterioration.

Experiment

The base material used in the experiment was DC 01=DIN EN 10130-9 bon deep-drawing steel with a thickness of 1 mm, treated by nitro-oxidationprocess in a fluidised layer The chemical composition of the material used isgiven in Table 1

waft of the fluidized medium during treatment was provided by gaseous nia A vapour of distilled water was supplied to the furnace chamber during oxi-dation The parameters of the treatment were as follows: nitridation at

The nitro-oxidation treatment according to the previously mentioned rameters influenced the base metal to the depth of about 260 lm in thickness,while the depth consisted of two zones (Fig 1)

pa-The previous experiments showed that the surface oxide layer caused someproblems when utilising GMAW and resistance welding methods [5] On thecontrary, the surface oxide and nitride layers are necessary for good corrosion-resistance properties; therefore, laser beam welding was applied The method iswidely used in the automotive industry for welding coated thin steel sheets [3]

Table 2 shows the welding parameters used

The nitro-oxidized material was welded by electron beam welding on a PZ

EZ ZH1 machine with various parameters at the First Welding Company, Inc.,Slovakia In order to exclude potential errors during the positioning and fixing

TABLE 1—Chemical composition of experimental steel [wt %].

C, Mn, P, S, Si, Al,

EN Code Max % Max % Max % Max % Max % Min %

DC 01 0.12 0.60 0.045 0.045 0.1 —

of the base materials, the bead-on-plate welds (hereinafter referred to only as

“welds”) were carried out on a piece of material The operating pressure in the

respectively Many welds were rejected due to porosity and the spatter observed.Regarding the previously mentioned findings, oscillation of the beam (forward-backward) was applied in the welding process Variable parameters in theresearch were as follows:

Tests were also carried out with multiple welding passes In order to obtain

a suitable porous-free weld, three passes of different parameters were required(Table 3) The first pass, considered a cleaning one, was removing the surfacelayer The second pass was the penetrating one, while the third one presentedjust a “cosmetic” modification of the surface

FIG 1—Microstructure of nitro-oxidized low-carbon steel DC 01

TABLE 2—Laser beam welding parameters.

Laser type Ferranti Photonics AF 8 CO2Protective gas Ar 99.996 % (18 l=min) Welding speed, v, mm=s 30, 40, 50, 60 Laser power, W 2000

condi-TABLE 3—Electron beam weld parameters carried out by three passes.

Welding Speed, mm=s

Focusing Current, mA

Frequency, Hz

Heat Input, J=mm

the layer becomes unstable The expansion of the vapour in vacuum is veryintense and the vapour’s dynamic pressure has a negative influence on the mol-ten weld met al As a result, excessive spatter was observed.

Figure 3 shows the surfaces, roots, and cross-sections of the welds as adependence on the EBW parameters of the specimens with the lowest spatterand lack of porosity

A close-up of the electron beam (EB) weld macrostructure is shown inFigure 4

The microstructure of the LBW joint (welding speed of 40 mm=s) is shown

in Fig 5, where particular areas of the weld joint can be seen—parent material,the heat affected zone (HAZ), and the weld metal

The microhardness measurements of the laser beam welds are shown inFigs 6 and 7 Table 4 summarizes the average value of the microhardness in therelevant section of the electron beam joint, while Fig 8 shows the microhard-ness progression through the previously mentioned weld joint

FIG 3—Surface, root, and cross section appearance of electron beam welds

18 JAI  STP 1532 ON 18TH IFHTSE CONGRESS

Visual inspection of laser beam welds proved incomplete penetration at thewelding speed of 60 mm=s When applying other welding speeds, full penetra-tion and regularly formed welds without defects were achieved The comparison

of the HAZ widths showed that the welding speed of 30 mm=s had the greatestheat affect As for the cross-sections, they looked similar at the welding speeds

of 40 and 50 mm=s, yet the weld root surface pattern exhibited some ities at the 50 mm=s welding speed The recommended optimum welding speedfor a thickness of 1 mm should thus range between 40 and 50 mm=s [7]

irregular-On the contrary, the welds produced by the EBW exhibited deficienciesregarding the weld bead surface regularity and excessive spatter [8] Since the

FIG 4—Macrostructure of a three-pass EB weld

FIG 5—Microstructure of a laser weld joint (weld metal- HAZ- parent material) of thematerials treated by nitro-oxidation

FIG 6—Microhardness as a dependence on the distance of the weld axis in the upperpart of a weld joint.

FIG 7—Microhardness as a dependence on the distance of the weld axis in the rootpart of the weld joint

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untreated materials did not exhibit a similar behaviour, we suppose that itmight be due to the presence of the surface layer The process of the surfacelayer deterioration during welding created vapour The expansion of the vapour

in vacuum is much more intense in comparison to the expansion at pheric pressure The vapour’s dynamic pressure can thus affect the molten weldmetal, while resulting in the excessive spatter

atmos-Multi-pass welding mode had, however, a positive influence on the weldbead surface regularity (see the comparison of the cross sections of samples 503and 502 in Fig 3) The shape of the weld cross sections was characterised by asurface concavity observed on all welds The direct comparison of the laser andelectron beam welds (Fig 2 versus Fig 3), showed a marked difference for thebenefits of laser beam welding, both in the weld bead dimensions and the shape

of the weld cross section, as well as in the amount of spatter and weld metalporosity

TABLE 4—Approximate values of microhardness of the electron beam weld.

Surface Middle Root Weld metal 152.7 159.1 185.7 Heat affected zone 161.3 134.5 181.6 Base material 154.8 124.8 180.2

FIG 8—Microhardness measurements of a three-pass EB weld

The width of the weld made by electron beam welding was double in parison to the laser beam welding, while the heat affected zone was even triplethe width This was probably caused by the lower cooling speed in vacuum dur-ing the electron beam welding, since the heat input of both welding methodswas almost the same (Table 1) The width of the heat affected zone of the elec-tron beam weld was very similar, regardless of the number of passes (Fig 3).The laser beam welding caused the nitride dissolution up to a distance of

com-1 mm from the weld joint boundary The microstructure of the weld metal sisted of acicular ferrite precipitating along the boundaries of columnar crystals(Fig 9)

con-Despite the anticipated problems regarding the occurrence of nitrides in thesurface layers of the materials, the structure did not exhibit any abnormalities

in phase composition

The microstructure of the high temperature HAZ consisted of polygonal rite The undesirable grain coarsening was not observed in this part of the weldjoint (Fig 10)

fer-Based on the microstructural analysis, we can assume that the existing

Having a partially deformed texture, the microstructure of the parent materialconsisted of polygonal ferrite and tertial cementite precipitated at the grainboundaries

The electron beam weld microstructure depended on the number of ing passes The microstructure of single pass welds consisting of polygonal fer-rite manifested a coarser grain structure (Fig 3; cross-section of sample 503).The presence of upper bainite was sporadically observed

weld-FIG 9—Microstructure of the weld metal of a laser weld

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The microstructure of the welds made by three passes was different, sinceevery subsequent pass had an influence on the microstructure of the previousone The microstructure of the upper part was formed by coarse acicular ferriteand polygonal ferrite Sporadically, grains of the upper bainite were alsoobserved (Fig 11) The ferritic matrix contained secondary precipitated car-bides and nitrides The structure of the weld root was very similar to the upperpart of the weld Significant differences in the weld microstructure wereobserved in the middle part of the weld This was due to the application of thethird welding pass resulting to the annealing of the previous, i.e., the secondpass The microstructure of the weld area consisted of fine polygonal ferritewith a slight heterogeneity of the grain size (Fig 12).

On the contrary, the microstructure of the high temperature heat affectedzone in the upper, middle, and root parts of the weld joint was very similar Itconsisted of polygonal ferrite with a fine and almost equal grain size No graincoarsening in this part of the weld was observed (Fig 13)

The microstructure of the base metal-heat affected zone transition for thethree monitored weld parts consisted of polygonal ferrite with precipitatednitrides in the ferritic matrix and at the grain boundary There was a significantdifference in the grain size of the heat affected zone on the one hand, and thebase metal structure on the other hand (Fig 14)

Microhardness measurements of the laser beam welds showed a maximumhardness of 360 HV 0.1 in the weld metal of the specimen welded by a 60 mm=swelding speed The specimen also exhibited the highest microhardness var-iance, which was probably due to the highest welding speed and thus, the high-est cooling rate This estimation can be supported by the microhardness

FIG 10—Microstructure of the base metal-heat affected zone transition of a laser weld

FIG 11—Microstructure of the weld metal in the upper part of the EB weld.

FIG 12—Microstructure of the weld metal in the middle part of the EB weld

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measurements of a specimen welded at the lowest welding speed (30 mm=s)which also had the lowest microhardness variance owing to the lowest coolingrate The differences in microhardness in the upper and root parts of the weldjoints were insignificant.

The microhardness measurements of electron beam welds showed the est microhardness of 120 HV 0.1 near the sample surface, and a bit highermicrohardness (172.6 HV 0.1) in the heat affected zone Another microhardnessdrop to approximately 150 HV 0.1 was observed in the base material near thesurface The middle section of the electron beam weld showed a stable micro-hardness of approximately 157 HV 0.1, continuously decreasing through theheat affected zone up to the base material to approximately 122 HV 0.1 Thehighest values of microhardness of approximately 180 HV 0.1 were observed inthe root section of the weld metal The fluctuations in microhardness were prob-ably due to the presence of upper bainite in the microstructure and the possibleplacement of an indent in this phase

low-Conclusions

The changes of the microstructure and properties of low-carbon deep-drawingsteel sheets after the nitro-oxidation process in a fluidised bed wereinvestigated

multi-phase composition of the sheets after their nitro-oxidation An undesirable

FIG 13—Microstructure of the EB weld metal-high temperature heat affected zonetransition

demonstration of a microstructural change in the analysed weld joints was notobserved The analysed microstructures of the weld joint areas were of a similarcharacter to those corresponding to the welding of structural steels.

The results of the electron beam welding of nitro-oxidation treated steelsheets were inconsistent When compared to laser beam welding, the width ofthe heat affected zone was three times wider and the weld width was double.Since the decrease of the heat affected zone width (important for minimis-ing the deterioration of the nitro-oxidation layer) was not shown, it is suggestedthat EB welding is not used when welding nitro-oxidized steels On the contrary,the laser beam welds were spatter-free and had a more suitable weld shape inthe cross-section Unlike the electron beam welding, the laser welds did notexhibit concavity of the weld surface

FIG 14—Microstructure of the base metal-heat affected zone transition of the EB weld

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[2] Lazar, R., Maroˆnek, M., and Doma´nkova´, M., “Low Carbon Steel Sheets Treated by Nitrooxidation Process,” Strojarstvo Extra, Vol 4, 2007, p 86.

[3] Varga, V and Vinˇa´sˇ, J., “Laser Beam Welding of Coated Thin Steel Sheets,” ceedings of Technolo´gia 2001: 7th International Conference, Bratislava, Sept 11–12,

Microstruc-[6] Maroˆnek, M., Ba´rta, J., Ba´rtova´, K., and Drˇı´mal, D., “Welding of Steel Sheets Treated by Nitrooxidation,” 16th International Conference on the Joining of Materi- als and 7th International Conference on Education in Welding ICEW-7, Tisvildeleje, Denmark May 10–13, 2011.

[7] Maroˆnek, M., Ba´rta, J., Lazar, R., Doma´nkova´, M., and Kovarˇı´kova´, I., “Laser Beam Welding of Steel Sheets Treated by Nitrooxidation,” 61st Annual Assembly and International Conference of the International Institute of Welding, Graz, Austria, July 6–11, 2008, pp 1–8.

[8] Maroˆnek, M., Ba´rta, J., Doma´nkova´, M., Ulrich, K., and Kolenicˇ, F., “Electron Beam Welding of Steel Sheets Treated by Nitro-Oxidation,” The 62nd Annual As- sembly and International Conference of the International Institute of Welding (IIW), Singapore, July 12–17, 2009, pp 1–11.

CRYOGENIC