of foundry engineering · csmx-4, inconel 738 at the temperatures of 700 and 900oc. however, it...

A R C H I V E S o f

F O U N D R Y E N G I N E E R I N G

Published quarterly as the organ of the Foundry Commission of the Polish Academy of Sciences

ISSN (1897-3310) Volume 10

Special Issue 3/2010

11 – 16

2/3

A R C H I V E S o f F O U N D R Y E N G I N E E R I N G V o l u m e 1 0 , S p e c i a l I s s u e 3 / 2 0 1 0 , 1 1 - 1 6 11

Microstructural changes in padding welds made from the 713C alloy after heat

treatment

M. B. Lachowicz

Wroclaw University of Technology, Institute of Materials Science and Applied Mechanics, ul. Smoluchowskiego 25, 50-370 Wroclaw, Poland

Corresponding author. E-mail address: [email protected]

Received 30.04.2010; accepted in revised form 30.05.2010

Abstract

The following paper presents the results of the metallographic research of the padding welds made from cast alloy Inconel 713C that had undergone heat treatment in four different time and temperature variants. Annealing was applied at temperatures 1150oC and 950oC. The results of the research showed a strong coalescence of precipitation of the intermetalic γ’ phase (Ni3Al), resulting in formation of subgrains. The coalescence of the γ’ particles ran simultaneously with a diffusive decomposition of carbides (NbTi)C, which caused the supersaturation of the γ’ phase with niobium, and the loss of its coherence with austenite γ. The phenomenon of the formation of subgrains intensified with the lengthening of the time of annealing. The analysis of the results showed that heating the alloy that had undergone pad welding to the temperatures used in the research is detrimental to the alloy because of the loss of the strengthening of the coherent γ’ phase and the increase of the brittleness.

Keywords: heat treatment, metallography, superalloy 713C, Ni3Al, NbC

1. Introduction Superalloy 713C is one of the oldest casting superalloys based

on nickel, strengthened by the intermetalic phase Ni3Al - γ’. The content of the intermetalic γ’phase in the microstructure exceeds 50% and its precipitations vary from 0.4 do 1μm [1.2]. The microstructure and mechanical properties of the In. 713C alloy constitute the frame of reference to a new generation of superalloys. The Inconel 713C alloy is characterized by one of the highest resistances to high-temperature corrosion in the group of superalloys. Its creep strength outstrips alloys such as Waspaloy, CSMX-4, Inconel 738 at the temperatures of 700 and 900oC. However, it manifests little plasticity (A≈6%), a feature which

magnifies its tendency to brittle cracking. During designing stages alloy 713C was not intended for welding purposes and it is still considered to be one of the least weldable superalloys based on nickel. In fact, a term unweldable alloy would be more appropriate [3].

Alloy 713C is most commonly worked just after casting. However, it can be submitted to heat treatment e.g. according to a pattern described by Somoza [1]. It consists in supersaturating at temperatures from 1150oC to 1180oC for 2 to 4 hours, and then age-hardening at a temperature of 950oC from 4 to 100h. In temperatures 1200oC and 1230oC, interoperational annealing might be applied to that alloy as well. [4]. Cooling can take place by means of a furnace, water or air.


2. Material and methodology of research

For the research purposes the Inconel 713C alloy was used, with the following chemical composition: Cr 12.8 %, Mo 4.15 %, Nb 1.73 %, Al 6.2 %, Ti 1.04 %, C 0.12 %, B 0.017 %, Zr 5 ppm, Fe 0.19 %, Ni – the rest. In these trials ceramic mould casting plates with the dimensions 120 x 70 x 5 mm were used.

The aim of the heat treatment was to homogenize the structure and eliminate the internal stresses in previously made padding and fusion welds. Yet another goal of the research was attempt to relieve the material stress and take measurements of the stresses in states after casting, welding and stress relieving. The outcome of the stress measurements done using an X-ray method was not included in this paper due to a vast range of tolerance, caused by a large grain in Inconel 713C alloy.

Four different variants of the heat treatment were applied in the research:

W1 – padding weld + base material - annealing in 1150oC for 1h and air cooling + 950oC for 4h and air cooling,

W2 – padding weld + base material - annealing in 1150oC for 4h and air cooling + 950oC for 16h and air cooling,

W3 – base material - annealing in 950oC for 30h and air cooling,

W4 – base material - annealing in 950oC for 30h and furnace cooling.

Both the samples and the furnace were heated gradually, with

a 15 min break at the temperature of 600oC. The samples had been cooled down from the temperature of 1150oC and then they were heated up again to the temperature of 950oC after the furnace had been cooled down to the temperature of 300oC.

In the described research both padding welds and the alloy in its initial state (after casting) were submitted to the heat treatment. The padding welds were made in Welding Department of Wroclaw University of Technology using the GTA method on the plates with the dimensions 70x30x5mm. The bars made out of Inconel 713C alloy with the diameter of about 2.5mm were used as filler metal. The padding welds were hand-made using variable current with the intensity of 70A and average voltage of 28V. The speed of welding was about 2mm/s (0.12m/min).

In the microscopic research both light and scanning microscopy were applied. Metallographic specimens were etched electrolytically using the ML3 method [5,6].

3. Research results

3.1. Microscopic research of the padding welds after the heat treatment W1 and W2

On the basis of observations of the microstructure in the

padding welds after W1 and W2 variant treatment, it was

concluded that homogenizing of the microstructure took place in both the padding weld and in the heat-affected zone (HAZ) in terms of size and quantity of γ’particles, through their growth and coagulation. The boundary zone of the padding weld - HAZ can be distinguished only thanks to the character change of the arrangement of the structure from the initial dendritic arrangement in the base material to smaller cellular-dendritic arrangement in the padding alloy. The microstructures in the base material, the HAZ and in the padding weld are presented in the fig.1.

Fig. 1. The microstructures of the HAZ and of an example of the padding weld after heat treatment W1. The visible change in the

padding weld of the dendritic arrangement from the initial to cellular-dendritic. The uniform colour of the γ-γ’ matrix in the

HAZ and in the padding weld. Etched using ML3

The microstructure of the base material, the padding weld and

HAZ observed through the light microscope were marked by the pale and uniform colour which is influenced by the size and number of γ’ particles. Continuous and complex precipitations of the phases, formed due to unbalanced decomposition of carbides, were observed during the research. This type of phases was described by means of stoichiometric formula in the paper [7] as Ni3(Al,Ti,Nb)C or Ni3(Al,Ti,Nb).

The post-carbide phases formed closed areas resembling grains, as well. Inside those subgrains a change in morphology and orientation of the particles of γ’ phase was observed, which mean that they can posses a different crystallographic texture. The areas in the padding weld with the visible post-carbide γ’phase, forming oblong precipitations and subgrains are shown in the fig. 2.


Fig. 2. The structure of the padding weld. Visible precipitations of the post-carbide γ’ phase forming subgrains in the uniform γ-γ’

matrix. Etched using ML3

The observations in which scanning microscopy was used, have shown that in the areas around the carbides γ’ particles can be found. They have got near-cube shapes and their size is smaller than γ’ phase’s distance form the carbides, shown in the fig. 3. It confirms a thesis that, γ’ particles undergo a dissolution in the areas around carbides in lower temperatures than particles in the areas distant from carbides – in axis of dendrites. The range of the decomposition temperature can start at just 1120oC or even in a slightly lower temperature which was demonstrated in the paper [7]. However, at the temperature of 1150oC the decomposition of the γ’ particle is already advanced.

Fig. 3. Dissolved and re-precipitated γ’ phase around the initial carbide in HAZ. Etched using ML3. SEM

In the areas distant from the recrystallized carbides in the padding welds, a distinct coagulation and coalescence of initial γ’ particles took place. The coalescence of the γ’ phase caused the increase of its size up to three times bigger that its initial size, which is up to 1 – 1.5μm. At the same time the distance between the precipitations increased as well. The coalescence of the γ’ particles happens in the first place between the carbides or the post-carbide precipitations and it is connected with the segregation of the chemical constitution in those areas. The elongated γ’ precipitation is shown in the fig. 4 while in the fig. 5 one can observe subgrains formed by the post-carbide areas.

Fig. 4. The precipitations of γ’ phase with the elongated shape that take place in the post-carbide areas and are surrounded again by

the precipitated γ’ phase. Etched using ML3. SEM

Fig. 5. Extended post-carbide precipitations of γ’phase in the padding weld. Etched using ML3. SEM


3.2. The analysis of the padding welds and the base material after W2 treatment

After the W2 variation treatment it was observed that there

was a significantly higher number of subgrains forming in the inter-carbide areas, both in the padding weld and in the base material. It shows that extending of the annealing time results in the intensification of the coalescence process of γ’ particles in the inter-carbide areas. On the basis of the observations of the microstructure it can be noted that newly formed subgrains resemble the grains occurring in the metallic materials worked plastically - fig. 6 and 7.

Fig. 6. The microstructure in the padding weld after the heat treatment W2. The visible, numerous subgrains are formed

through the coalescence of γ’ particles in the inter-carbide areas and single pale γ’ particles. Etched using ML3

Fig. 7. γ’ phase forming subgrains, precipitated and connecting initial carbides in the base material. There is a visible pale, initial γ’ phase and partially recrystallized γ’ particles in the matrix. .

Etched using ML3

While observing the microstructure of the padding welds and the base material using the scanning microscopy, it was concluded that the extension of the time of the heat treatment caused both a higher homogenizing of the microstructure and dissolution of the bigger number of γ’ particles. In the microstructure there are simultaneously γ’ particles 0.5μm big that have a cubical shape with slightly rounded corners, and between them there are bigger particles with a size approximately 1.5 do 3μm. The bigger γ’ particles are marked by oval shape, however, they can also occur as regular parallelograms with a base of a square or a rectangle. The microstructure in the base material is shown in the fig. 8.

Fig. 8. An individual bigger γ’ precipitations and post-carbide γ’ precipitations forming subgrains. The carbides are visible as pale precipitations against the background of the matrix γ-γ’. Etched

using ML3. SEM, BSE The occurrence of γ’ phase with such a diverse shape can be

explained by dissolution of all γ’ particles existing in the padding weld during the annealing at the temperature of 1150oC, and subsequently by their precipitation, coagulation and coalescence. It is also probable that, the bigger particles did not manage to dissolve but they underwent coalescence which is a result of a higher supersaturation of niobium of γ’ phase in the areas distant from the carbides. In his research [1] Somoza conducted supersaturation at the temperature of 1170oC, and age-hardening at the temperature of 950oC for 750min (12.5h) which allowed him to conclude that the size of the particles has reduced from 0.5μm (before supersaturation) to 0.25μm. Whereas the lengthening of the annealing time to 2000 min led to the re-growth of the particles to 0.5μm. In his research, however, Somoza does not state whether the oblong γ’ precipitations occurred in the post-carbide areas.

On the basis of the conducted analyses of the superficial distribution of the elements, it should be concluded that bigger γ’ particles, similarly to oblong γ’ precipitations are richer in niobium. The higher content of niobium results from the decomposition of the re-crystallized carbides in the padding weld.

In the pictures of the microstructure, taken with the use of the contrast material BE and shown in the fig. 8 and 9 one can


observe that the γ’ precipitation forms around the carbide and it joins carbide-forming elements which is mostly niobium. The superficial elements distribution in bigger, oval precipitations is presented in the fig. 9. The increase of the niobium content rises the temperature of the decomposition of the γ’ particle [5-7]. It encourages the formation of subgrains through the coalescence of γ’ phase in the post-carbide areas and between the carbides as well. Additionally, newly formed γ’ particles can be precipitated on the initial carbides and then a diffusive decomposition of a carbide, and the supersaturation of the γ’ phase with carbide-forming elements can take place.

Fig. 9. The superficial distribution of the content of niobium - B. A visible increase in the contents of niobium in the post-carbide

precipitations. A- the microstructure with a visible γ’ phase surrounding lighter carbides NbC

3.3. The microscopic research after W3 and W4 heat treatment variants

Lengthening of the annealing time in the material, both before

and after welding increases the amount of subgrains forming in the inter-carbide areas as well. It is proven by the examination of the sample of the base material put to the annealing only at the temperature of 950oC for 30h – W3. The amount of the subgrains formed by incoherent γ’ precipitations was significantly higher, which is shown in the fig.10. On the grain-boundary of the γ-γ’ matrix a continuous envelope from the incoherent γ’ phase occurred as well. The morphology of γ’ particles of the grains of the matrix was basically unchanged; the matrix has undergone

only a slight rounding of its corners. At the applied temperature the dissolution of the γ’ particles did not take place. It should be concluded then, that application of that treatment before welding, for instance in order to remove internal stresses or to homogenize the structure is purposeless. Annealing in those parameters will raise the amount of the incoherent phases and will cause the increase the brittleness of the alloy.

Fig. 10. The microstructure of the sample of the base material after heat treatment W3. Visible numerous subgrains are formed

through incoherent γ’ phase precipitations in the inter-carbide areas. Etched using ML3

Heat treatment reduced the hardness of the padding welds

from 450±8 HV1 to 396±6 HV1. The hardness of the base material after the heat treatment W3 was 369 ± 5 HV1, whereas its initial hardness was at approximate level of 394± 6 HV1.

The change of hardness of the padding welds was not strictly confined to the applied variation of the treatment. The hardness of the padding welds, however, has decreased significantly in relation to the value just after pad welding which is connected to the transformation and the growth of γ’ nanoparticles after the treatment and the reduction of the precipitating strengthening of the padding weld. The higher level of the hardness of the padding weld than of the base material confirms the theory that there is still a strengthening in the microstructure of the padding weld caused by the supersaturation of the matrix, caused most probably by the carbide-forming elements. The sample of the base material which had undergone the heat treatment W4 did not manifest noteworthy differences in comparison to the sample after heat treatment W4.

4. Conclusions

The coagulation and the coalescence of γ’ particles and the formation of oblong precipitations and subgrains led to the loss of their coherence and may cause the increase of the brittleness of the alloy. Most probably it is be caused by the formation of additional misfit dislocations and blocking of the dislocation in the interfacial boundaries, which may lead to stress fractures.


Because of the changes in the microstructure of the γ’ phase transformation, heating the alloy to the temperatures applied in the study should be avoided. The temperatures of the heat treatment before and after welding must be lower than 950oC, in order not to cause detrimental changes of the microstructure but just to reduce the stresses.

Acknowledgements

The research was conducted as part of a programme - Fellowship co-financed by European Union within European Social Fund

References

[1] A. Somoza, G. Santos, A. Ges, R. Versaci, F. Plazaola, Age-hardening and precipitation phenomena in the Inconel 713C superalloy studied by means of positron lifetime spectroscopy, Phys. Stat. Sol. 174, (1999) 189-198.

[2] F. Zupanič, T. Bončina, A. Križman, F.D. Tichelaar, Structure of continuously cast Ni-based superalloy Inconel 713C, J Alloys Compo 329, (2001) 290-297.

[3] M.B Henderson, D. Arrell, M. Heobel, R. Larsson, G. Marchant, Nickiel-Based Superalloy Welding Practices for Industrial Gas Turbine Applications, www.msm.cam.ac.uk/ phasetrans/2002/papers/APNickelWeldv2.

[4] B. Mikułowski; Heat-resisting and creep-resisting alloys – superalloys; AGH, Kraków, 1997 (in Polish).

[5] M. Lachowicz, W. Dudziński, K. Haimann, M. Podrez-Radziszewska, Microstructure transformations and cracking in the matrix of γ-γ’ superalloy Inconel 713C melted with electron beam, Mat Sci Eng A 479 (2008) 269-276.

[6] M. Lachowicz, W. Dudziński, M. Podrez-Radziszewska, TEM observation of the heat-affected zone in electron beam welded superalloy Inconel 713C, Mater Charact 59 (2008) 560-566.

[7] M. Lachowicz, W. Dudziński, Non-equilibrium decomposition of MC carbides in superalloy Inconel 713C melted with welding techniques, Archiv Metall Mater, vol.55, No 1 (2010) (in the press).

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