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Vol. 8. Issue 2.
Pages 1636-1644 (April 2019)
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Vol. 8. Issue 2.
Pages 1636-1644 (April 2019)
Original Article
DOI: 10.1016/j.jmrt.2018.11.009
Open Access
Study of the high temperature oxidation and Kirkendall porosity in dissimilar welding joints between FE-CR-AL alloy and stainless steel AISI 310 after isothermal heat treatment at 1150°C in air
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André de Albuquerque Vicentea, Joao Roberto Sartori Morenob,
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joaosartori@utfpr.edu.br

Corresponding author.
, Denise Crocce Romano Espinosaa, Tiago Felipe de Abreu Santosc, Jorge Alberto Soares Tenórioa
a DSão Paulo, Rua do Lago, 250, Cidade Universitária, São Paulo, SP, Brazil
b Mechanical Engineering Department, Universidade Tecnológica Federal do Paraná, Alberto Carazzai, 1690, Cornélio Procópio, PR, Brazil
c Department of Mechanical Engineering, Universidade Federal de Pernambuco, Av. da Arquitetura, s/n, Cidade Universitária, Recife, PE, Brazil
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Tables (3)
Table 1. Welding parameters of test specimens.
Table 2. Chemical compositions of the base metals and the welding consumable (wt%).
Table 3. Chemical compositions of base metals, welding consumables and all weld metals, before and after heat treatment (wt%).
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Abstract

Dissimilar welded joints of Fe-Cr-Al alloy (23Cr-7Al) and AISI 310 austenitic stainless steel (25Cr-20Ni) were studied and characterized, before and after the isothermal heat treatment (IHT) of oxidation at 1150°C for 936h in air. The Fe-Cr-Al alloy was used as the welding consumable and the welding process used was gas-shielded tungsten arc welding (GTAW) with currents of 60, 80 and 130amps, and welding speed 1, 1.1 and 1.3mm/s. In the heat treated specimens, pore formation was observed as well as significant variations of Al contents in all-welded metals and formation of protective layer of alumina-Al2O3 on the oxidized surfaces particularly in the sample 1, welded at 60amps and welding speed of 1mm/s, that showed higher pore index and a greater deflection in the decrease of the Al content. The formation of Kirkendall pores occur due to the rapid diffusion of the aluminum from the all weld metals to the AISI 310 austenitic stainless steel, as was detected by the sample 1 mainly and especially by identifying the chromium content of the dissimilar alloy and the stainless steel.

Keywords:
GTAW
Kirkendall porosity
Alloys Fe-Cr-Al
AISI 310
Full Text
1Introduction

Certain authors [1] optimized a body by replacing some steel parts with aluminum alloys, and the crash results show that there was little difference in the overall deformed shapes of the two designs, and most importantly, the optimal multi-material body achieved a weight reduction of 13%, with the material cost increasing by 7%. However, some disadvantages still exist in the multi-material approach, and the most notably one is the challenge of joining dissimilar materials.

Due to their different thermophysical properties and diffusion these alloys [2] and diffusion studies can be divided into pure metals, homogeneous alloys and heterogeneous alloys. In the study for heterogeneous alloys, the understanding of the mechanisms of diffusion can be in a macroscopic scale, being evaluated the variations of concentration of the solutes [3].

As the diffusion effect [4], must occur at the grain boundary the effective diffusion coefficient is related to the isothermal treatment conditions, and the specific phase and its grain size are important to find the influence of these factors on the deposition. Therefore, the size of each phase is certainly the type of dissimilar alloys and the grain boundaries and grain growth during the welding procedure.

For a system of alloys with elements A and B with different diffusion coefficients, compared to the pure metals, the mechanism of diffusion by vacancies requires less energy of diffusion. Kirkendall porosity was demonstrated in experiments with Cu-Zn alloys that zinc diffused with a higher velocity than copper and that the formation of voids in the Kirkendall experiment in terms of preceding analyses of vacancy diffusion. In particular, use is made of the viewpoint that the presence of voids is related to a super-concentration of vacancies in some region of the diffusion zone [5].

Therefore, dissimilar welding becomes an interesting area, and a large amount of research work is being carried out on welding different combinations, such as aluminum to steel, aluminum to titanium and others.

Have a lot of work on optimization of welding aluminum to steel [6,7], where the relationship between welding parameters and mechanical property was reported, and the microstructure of base metals was observed.

Various types of ferritic materials have been developed to have resistance to corrosion and oxidation at high temperatures. They are characterized by having low carbon content and high chromium content and as chromium is a ferrite stabilizer it cannot be hardened by heat treatment.

The excellent oxidation resistance of Fe-Cr-Al alloys as seen in Fig. 1[8], is due to the formation of Al2O3 in the zone at 1100°C is bounded by a broken line where the zones of Fe3Al and Cr-enriched α′ phase at 475°C are bounded by a dotted line and a solid line, respectively. Based on Fig. 1 of the Fe-Cr-Al ternary diagram, we can also observe that in the interaction of this alloy with high chromium austenitic steel, such as AISI 310 steel, the predominant oxides after treatments around 1100°C are Al2O3 and Cr2O3 in function of the high and facilitated diffusivity of the oxygen.

Fig. 1.

Fe-Cr-Al ternary diagram [8].

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These alloys are used in the manufacture of electrical resistances, and are more cost effective than austenitic stainless steel, making them very expensive in some applications. When higher mechanical strength is required, austenitic steels, or in some cases chromium-nickel alloys, should be used [9].

Although the temperature rise accelerates the layer loss process, even when operating at temperatures up to 1400°C (from 100 to 200°C below melting point), an alloy with 15–30% chromium and at least 5% of aluminum in weight, loses only 0.5–1.0g/m2/h, which is less than the losses recorded in heat resistant nickel chrome steels.

Authors define that for an alloy containing 25% chromium, 5% aluminum in weight, after exposure at a temperature of 1200°C for 240h resulted in 98.72% Al2O3, 0.80% Fe3O4 and 0.68% Cr2O3[10,11].

Even after prolonged periods of exposure, the chromium content of Fe-Cr-Al alloys remains the formation of the passivation layer (oxide) depends on the aluminum content of the alloy. In alloys with low aluminum content the drop may reach 60%, while in alloys with high content the average loss of aluminum is very low. The formation of the oxide layer at high temperature occurs according to the following reactions:

2(Al, Cr, Fe)+4.5O2Al2O3+Cr2O3+Fe2O3
Cr2O3+2AlAl2O3+2Cr
Fe2O3+2AlAl2O3+2Fe
At temperatures up to 700°C the scale consists of a solid solution of three oxides above and that this process is controlled by Eq. (1). When the temperature is raised above 700°C, the aluminothermic reactions represented by Eqs. (2) and (3) occur simultaneously with the reaction represented by Eq. (1).

In this way the thickness of the oxide film remains in growth and both the Cr2O3 chromite and the Fe2O3 hematite are reduced to form Al2O3 alumina.

The Kirkendall effect manifests itself by the movement of marker planes in a diffusion couple system. It is the result of a net mass flow, accompanied by a vacancy flow in the opposite direction [12].

While there are numerous applications for dissimilar joints between steel/Al and alloy, there remain several open questions regarding the fundamental role of intermetallic phases in joining of Cr-Al alloys to steel. These are: (a) the chemical potentials, (b) the nucleation conditions at the beginning of the interdiffusion process, and (c) the mobilities of the constituent elements.

Understanding of the fundamental reactions governing the phase formation and evolution at the interface between Fe/steel/Al; especially at temperatures below the melting point of Al; Exploration of the mechanisms how intermetallic reaction layers influence the mechanical properties of dissimilar joints between steel and Al alloys [13].

Although the specific properties that define the Fe and Al alloys are more difficult by the great difference in the melting temperatures of these metals, there may be the formation of brittle intermetallic particles which alter these properties. The Al-Fe phase diagram shows a high solubility of Fe for Al and three Al-rich intermetallic phases: ζ (Al2Fe), η (Al5Fe2) and θ (Al13Fe4) [14].

Diffusional growth of intermetallic compounds and elucidation of the Kirkendall effect accompanying solid-state reactions are subjects of considerable complexity, having broad applicability in materials science and engineering [15].

2Experimental procedure

Three specimens with different welding parameters, corresponding to the parts used in the furnaces, were welded. An AISI 310 stainless steel round bar of dimensions Ø 9.5mm×100mm in length and wires of an alloy Fe-Cr-Al (alloy Kanthal A1) of sizes Ø 4.0mm×100mm in length were used. Fig. 2 shows the dimensional details of this specimen.

Fig. 2.

Dimensional of the cross section in mm and details of the cylindrical specimen.

(0.11MB).

The groove welding was performed with a 45° angle that allowed the access of the end of the wire. The GTAW (also known as TIG) process was used. The protection gas used was pure Argon and the welding source, constant current. Current (A) and welding speed (mm/s) were monitored during welding for further calculation of welding energy.

The welding of the specimens was performed in the flat position. Table 1 lists the welding parameters of the specimens.

Table 1.

Welding parameters of test specimens.

Specimen  Polarity  Current (A)  Tension (V)  Welding speed (mm/s)  Heat input (kJ/mm) 
CP1  CC-  60.0  9.8  1.0  0.59 
CP2  CC-  80.0  10.6  1.1  0.71 
CP3  CC-  130.0  10.3  1.3  1.03 

Chemical analyses of the all weld metals in all the test specimens were carried out using an optical emission spectrometer, according to ASTM E 1086-08 [16]. Table 2 shows the chemical compositions of base metals and consumable used, according to the manufacturer's quality certificates.

Table 2.

Chemical compositions of the base metals and the welding consumable (wt%).

Component  Fe  Cr  Al  Si  Ni  Zr  Mo  Mn 
AISI 310 – Ø 9.5mm  52.18  25.01  0.00  0.06  20.11  0.00  0.35  1.93 
Kanthal A1 – Ø 4.0mm  68.96  23.23  7.30  0.33  0.08  0.09  0.00  0.00 
Kanthal A1 – Ø 2.0mm  70.12  21.41  7.08  0.30  0.08  0.08  0.00  0.00 

An isothermal heat treatment (IHT) was carried out at 1150°C for 39 days, totaling 936h that was carried out in a muffle-type oven, in several samples. For the preparation of samples for metallography, the specimens were cut in the region of the welded joint.

Subsequently, the samples were hot mounted in resin bakelite. The conventional manual grinding was used using sandpapers (100, 240, 320, 400, 600 and 1000 mesh) in order to standardize the surface finish of the samples. Afterwards, the samples were polished with a 9, 3 and 1μm diamond abrasive paste in this sequence.

The electrolytic attack with 10% oxalic acid was used, with an approximate time of 90s and a voltage of 5V. The equipment used was of the brand STRUERS, Polectrol model.

This allowed the microstructural characterization of the samples by optical microscopy, as well as a better SEM observation of the phases present in the oxidized layers. For oxidized samples was studied using X-ray diffraction. For this purpose, a RIGAKU X-ray diffractometer, MiniFlex model was used.

3Results and discussion

Base metals and welded joints Fe-Cr-Al alloy (Kanthal A1) were characterized before and after the IHT at 1150°C as shown in the micrographs of Fig. 3. However in Fig. 4 shows the micrographs of the base metals AISI 310 austenitic stainless steel as rolled, with an austenitic structure defined before and after the IHT.

Fig. 3.

Fe-Cr-Al base metal (Kanthal A1): (a) before IHT and (b) after IHT.

(0.36MB).
Fig. 4.

Base metal AISI 310 stainless steel sheet: (a) before IHT and (b) after IHT.

(0.39MB).

However, a macrograph study had to be performed to observe the levels of porosity distribution in the regions of the weld. With this, Fig. 5 shows it is macrographic analysis of welded in three different specimens CP1, CP2 and CP3.

Fig. 5.

Macrographs of the welded specimens: (a) CP1, (b) CP2 and (c) CP3.

(0.7MB).

But with the detailed analyses of Figs. 3–5, was observed the absence of porosities before the heat treatment.

On the other hand it was observed that in both base metals before the IHT we had more elongated grains. This microstructure is typical of thermomechanically shaped base metals, as in the case of the alloy Fe-Cr-Al (Kanthal A1) which was drawn, and the AISI 310 austenitic stainless steel which was rolled.

However we also noted that IHT, was characterized by the recrystallization of the elongated grains to equiaxial disposition. The chosen IHT temperature at 1150°C is approximately the solubility temperature of the stainless steel 310 which favors metallic diffusion.

Fig. 6 shows the micrographs of Kanthal A1/310 dissimilar joints, welded prior to heat treatment, where none of the pores are observed in base metals or metals deposited prior to IHT for the 3 specimens.

Fig. 6.

Micrographic characterization of welded joints before IHT. Dissimilar joints (Kanthal A1/310) details: (a) CP1, (b) CP2 and (c) CP3.

(1.05MB).

As the thermal input due to high currents (80–130A), there was a greater dissolution of the aluminum during the welding, occurring an instability in the formation of Al2O3 and consequently a decrease of pores in the weld metal as in shown Fig. 5b and c, respectively.

Fig. 7 shows the micrographic analysis of the region near the fusion line between the Fe-Cr-Al base metal (Kanthal A1) and the all weld metal on the welded joint of CP1, before and after IHT.

Fig. 7.

Region near the fusion line between Fe-Cr-Al base metal (Kanthal A1) and the weld metal of CP1 before and after IHT: (a) before IHT and (b) after IHT.

(0.38MB).

A traditional approach describes the pores nucleation in terms of classical nucleation from the metal core supersaturated with vacancies produced by Kirkendall effect. However, this type of homogeneous nucleation was shown to be unfeasible even in the bulk interdiffusion zone [17]. This makes the nucleation of pores in nanoparticles, which often lack the crystalline defects (and, thus, heterogeneous nucleation sites), even more problematic.

In turn in Fig. 8, it is observed the formation of a large amount of pores in the all weld metals, specifically in the regions near the fusion lines with the base metals AISI 310, in the three welded specimens after the heat treatment.

Fig. 8.

Micrographic characterization of welded joints after IHT: (a) CP1, (b) CP2 and (c) CP3.

(0.62MB).

The precipitations led to formation of precipitate base on Fe-Al described during rapid heating and isothermal annealing at 800°C as well as during a slower continuous heating to 900°C. During slower heating, the formation of Fe2Al5 and FeAl2 intermetallics starts below the melting point of aluminum and when the heating rate is high, intermetallics are created after melting of aluminum.

Fig. 9 shows the micrographic analysis of the region near the fusion line between the AISI 310 austenitic stainless steel base metal and the all weld metal of the CP1 joint before and after IHT.

Fig. 9.

Region near the fusion line between the AISI 310 austenitic stainless steel base metal and the all weld metal of the CP1 joint before and IHT: (a) before IHT and (b) after IHT.

(0.29MB).

It is known that in joints with pure Al, the mechanical properties are related to the formation of Kirkendall-porosity at the reaction layer/Al interface at where the Kirkendall is associate the existence of vacancy diffusion in the vast majority of metallic materials [18].

Precipitates are observed in the grain boundaries and porosities, both in the all weld metal and in the base metal near the fusion line, with an increase in grain size was also observed.

The void formation in these systems is the result of pure Kirkendall-porosity formation, because it is caused mainly by the inequality of the intrinsic atomic fluxes and other effects (e.g. stresses), inevitably present during nanoshell formations in solid state reactions (oxides, sulphides), can be less important or can be neglected [19].

The formation of intermetallic reaction layers for interdiffusion between steel and aluminum alloy (Al–5%) is characterized by the reactions Solid/solid, solid/semi-solid and solid/liquid diffusion couples were produced about 600 and 675°C.

The total width of the reaction layer is governed mainly by the parabolic diffusion-controlled growth of the phase (Al5Fe2), which exhibits orientation-dependent growth kinetics.

An important precondition for the successful wetting between steel and aluminum is the initial removal of the oxide layer from the aluminum surfasse, which in turn will migrate to the Kanthal layer in the form of Al2O3, which occurs by diffusion toward the vacancies, producing the porosities [20].

The chemical compositions on the surfaces of the all weld metals of the three specimens, before and after heat treatment, were analyzed and the results are presented in Table 3.

Table 3.

Chemical compositions of base metals, welding consumables and all weld metals, before and after heat treatment (wt%).

Component  Fe (±0.05)  Cr (±0.07)  Al (±0.08)  Si (±0.01)  Ni (±0.05) 
AISI 310 – Ø 9.5mm  52.18  25.01  0.00  0.06  20.11 
Kanthal A1 – Ø 4.0mm  68.96  23.23  7.30  0.33  0.08 
Kanthal A1 – Ø 2.0mm  70.12  21.41  7.08  0.30  0.08 
All weld metal CP1 before HT  63.69  23.09  4.46  0.32  7.93 
All weld metal CP1 after HT  63.24  23.25  4.41  0.31  8.29 
All weld metal CP2 before HT  65.99  22.66  5.40  0.30  5.22 
All weld metal CP2 after HT  63.56  23.62  3.77  0.31  8.21 
All weld metal CP3 before HT  65.31  22.99  5.04  0.30  6.09 
All weld metal CP3 after HT  64.35  23.31  3.89  0.32  7.27 

Analyzing Table 3, it was observed the decrease of the aluminum content in the all weld metal of the specimens after the IHT, independent of the welding energy used in the procedure, despite being on the margin of error 0.08%.

In this way, the analysis of oxides formed on the surfaces of the all weld metals of the specimens became interesting.

Fig. 10 shows the X-ray diffraction of the all weld metal surface of CP1 (Fe-Cr-Al alloy) after isothermal treatment at 1150°C for 936h in air, because it was CP1 sample that showed greater intensity of pores in the weld metal.

Fig. 10.

X-ray diffraction patterns of Fe-Cr-Al alloy.

(0.08MB).

The diffractogram showed well defined peaks of ferrite and the presence of the protective oxides Al2O3 and Cr2O3, with the protective oxide layer formed mainly of Al2O3, because in Fe-Cr-Al systems the small peaks of impurities, maybe pores, are related in peaks in accordance of Fe5C2 and Al2O3, as shown in Fig. 10.

X-ray diffraction was very useful technique for the analysis of deposits on metal surfaces, in the deposition metal only, as in the case of oxide layers formed at high temperatures [21].

The oxide film forming at temperatures below 700°C is made up of iron oxide, chromium and aluminum roughly in the same proportion as the alloying elements, however the extent of the scale was small. The film formed in aluminum-rich alloys subjected to temperatures greater than 700°C is light gray in color and is normally composed of pure oxide.

The migration of inert Kirkendall markers as a result of unequal mobilities of the components during solid-state interdiffusion in a binary system can be rationalized using the Kirkendall velocity diagram.

In a diffusion-controlled interaction the Kirkendall plane (identified by inert particles placed at the initial contact surface of a reaction couple), need not be unique. Multiple planes can be developed but, on the other hand, the Kirkendall plane after interaction can be unstable, i.e. markers can get dispersed into the diffusion.

In fact, the average diffusivity of the elements does not clarify the diffusivity of the species separately, that is, at the initial stage, the species diffusivity may be the same, but it is rearranged by the migration of the elements in the oxide form.

X-ray diffraction was applied [22] to the oxide surface through the semi-quantitative analysis at the metal / oxide interface.

Fig. 11 shows the micrograph, the qualitative chemical analysis and the characterization of the oxides formed on the surface of the all weld metal of CP1 (Fe-Cr-Al alloy) after isothermal treatment at 1150°C for 936h in the air.

Fig. 11.

Diffractogram of the oxidized surface of the all weld metal of CP1 after isothermal treatment at 1150°C for 936h in the air.

(0.1MB).

According to the data acquired by the diffractogram and the semi-quantitative analysis obtained by scanning electron microscopy with coupled EDS, of Fig. 12, it was possible to conclude that the protective oxide layer formed on the surface of the all weld metal of the Fe-Cr-Al alloy after an IHT at 1150°C for 936h in air, is formed mainly of Al2O3.

Fig. 12.

Diffractogram of the oxidized surface of the all weld metal of CP1 after IHT at 1150°C for 936h in the air.

(0.1MB).

According to the results of Figs. 11 and 12 we observed clearly that after the IHT the formation of Al2O3 and Cr2O3 chromium oxides was intensified on the surface of the weld metal and inside the pores.

4Conclusions

The formation of pores in the all weld metals occurs after IHT, since in the state as weld they were not observed. The aluminum content varied significantly in the all weld metals after the heat treatment and the layers of protective oxides observed on the all weld metals surfaces are mainly of Al2O3.

The obtained results suggest that the impurities observed in welded joints after heat treatment are probably Kirkendall pores, because in Fe-Cr-Al systems the small peaks of impurities related by the diffractogram in accordance with Fe5C2 and Al2O3 peaks.

The formation of Kirkendall pores occur due to the rapid diffusion of the aluminum from the all weld metals to the AISI 310 austenitic stainless steel, as was detected by the CP1 sample mainly and especially by identifying the chromium content of the dissimilar alloy and the stainless steel.

Conflicts of interest

The authors declare no conflicts of interest.

Acknowledgements

To SMT-SANDVIK Materials Technology, for believing in this project and for donating base metals and welding consumables. The UTFPR – Federal Technological University of Paraná for the technological support of the researches and USP – Escola Politécnica São Paulo, UFPE – Universidade Federal de Pernambuco, – Recife and the technicians who assisted in the analysis.

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Copyright © 2018. Brazilian Metallurgical, Materials and Mining Association
Journal of Materials Research and Technology

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