DEFECT FORMATION OF GRAY IRON CASTING Abstract

49International Journal of Metalcasting/Fall 09

DEFECT FORMATION OF GRAY IRON CASTING

A. Diószegi, L. Elmquist, J. Orlenius and I. Dugic

Jönköping University, Jönköping, Sweden

Copyright © 2009 American Foundry Society

Abstract

Cast iron is one of the oldest technical alloys (engineered

cast materials) used for creating objects. From the very

beginning of casting, foundrymen were fighting to avoid

casting defects. In the beginning a successfully produced

casting was associated with witchcraft. Cast component

producers suffer substantial yearly expenses due to

rejecting or repairing castings. This present work will

summarize research efforts to understand the formation

mechanisms of defects, performed in collaboration with

Swedish foundries during the last few years. It will focus

on defects specific to the casting of gray iron components

and the studied defects: gas porosity, shrinkage porosity

and metal expansion penetration. Novell experimental

set up has been developed or existing methods have been

improved to study defect formation mechanisms. Today

we can realize that casting without defects is possible

only by approaching the defect formation mechanism with

multidisciplinary science.

Keywords: cast iron, defects, gas porosity, shrinkage, metal

expansion penetration, pinholes

Introduction

The research in this report has been performed within a

collaborative project between Swedish foundries under the

umbrella of Casting Innovation Center and a Swedish ini-

tiative to promote the cast metal research. The performed

work focused on understanding the formation mechanism

of common casting defects such as gas porosity, shrinkage

porosity and metal expansion penetration in gray cast iron

components appearing at the collaborating foundries and

taking action to eliminate or minimize their impact on the

produced cast quality. Results from the research work have

been reported in journals and conferences ending up with

two PhD theses1,3 and a Licentiate thesis2. The Licentiate

thesis is a Swedish academic degree corresponding roughly

60 % of a PhD degree.

Gas Porosity

Gas porosity defects characterized in the literature as pin-

holes,4,5,6 are not only aesthetic problems but also counter-

act the functionality of complex shaped cast components.

These types of defects are revealed for the first time during

the machining, process thus causing additional expenses.

Two types of pinhole defects (hydrogen and nitrogen) are

similar in appearance but different in formation mechanisms

and have been reported in literature. In addition, pinholes are

also linked to carbon monoxide.4,5,7 The formation mecha-

nism of pinholes have been linked to bubble formation from

gases dissolved in liquid iron. Hydrogen and nitrogen are

soluble in liquid iron and are dependent upon the raw mate-

rials used, the melting method or the exposure of the metal

to a gaseous environment during mold filling. The solubility

of hydrogen in gray iron8,9 is approximately 8 ppm in liquid

state and drops to 4 ppm during solidification. The solubility

levels of nitrogen are reported below 150 ppm,9,10,11 but po-

rosity defect formation has been reported when the nitrogen

content is 110 – 150 ppm.12,13

Gas Concentration in Gray Cast Iron Melted in

Cupola And Induction Furnaces

The influence of the melting process on the hydrogen and ni-

trogen content in grey cast iron was investigated at two dif-

ferent foundries. Both foundries produce iron using induc-

tion furnaces, one also uses a cupola furnace. Measurements

of gas concentrations were performed at different places in

the process chain. In Foundry 1 the measurements were per-

formed directly in the induction furnace, while in Foundry

2, the liquid iron both from the induction furnace and the

cupola furnace, was transferred in a transport ladle of 3

tones before performing the measurement. Accordingly, the

chemical composition varied somewhat, but the lower and

upper limits are shown in Table 1. The hydrogen concen-

tration was determined by an online sensing system based

on thermal conductivity measurement of nitrogen gas cir-

culated between the measuring unit and the liquid iron until

equilibrium is achieved between the circulated nitrogen and

the hydrogen dissolved in the liquid. The average concentra-

tion of hydrogen in iron from induction furnaces were 2.1

and 2.7 ppm, while iron from cupola furnace has a concen-

tration of 1.4 ppm. The differences between the induction

furnaces are explained by the preheating of the charge when

the lower value is obtained. When preheating the charge ma-

terial, the amount of moisture is reduced contributing to a

lower hydrogen concentration. The cupola furnace produces

50 International Journal of Metalcasting/Fall 09

iron with an even lower concentration of hydrogen based

on the same principal of effective preheating of the charge

material. The nitrogen concentration was determined with

optical emission spectroscopy together with the chemical

composition from coin shaped samples collected in resin

bounded sample molds. The measured concentration of ni-

trogen by this method includes both dissolved and bounded

nitrogen. Cupola iron contains the highest concentration of

nitrogen, 105 ppm. Differences between the induction melt-

ing methods are also observed, 75 and 95 ppm respectively.

The highest nitrogen concentration from the cupola furnace

is attributed to the high amount of steel scrap the coke used

as fuel and the air blown in to the cupola. Under the investi-

gated conditions, neither the induction nor the cupola melt-

ing method results in hydrogen or nitrogen concentrations

close to the solubility limit reported in the literature except

for the nitrogen in cupola iron.

Gas Absorption in Gray Cast Iron

During Mould Filling

A special mould was developed to investigate the absorp-

tion of hydrogen and nitrogen in the liquid iron during the

mould filling. The scope of the investigation was to measure

the gas concentration in a pouring ladle before casting, then

cast the liquid into a special deigned gating system allow-

ing a large but controlled interaction between the liquid and

the surrounding mould components, and finally collect the

liquid in an open feeder where the hydrogen and nitrogen

concentration can be measured again. The moulds were con-

structed in such a way that different parameters influencing

the absorption process can be tested. The first parameter var-

ied was the gating procedure in order to create a different

level of turbulence in the gating system. The standard filling

concept is shown in Fig 1. The liquid metal is poured in the

mould through a pouring cup placed over a downsprue fol-

lowed by an ascending flow in a zigzag shaped channel up to

the open feeder. To provoke a more turbulent flow compared

to the standard filling concept, the liquid was poured into the

mould through the open feeder allowing for an uncontrolled

flow through the zigzag channel and finally ascending back-

ward up in the downsprue and finally filling the pouring cup.

The third variation of the filling sequence (Fig. 2) was ar-

ranged by introducing two pressed ceramic filters with the

intention to reduce the flow velocity. For the visual obser-

vation of the mould filling, the front side of the mould was

covered with thermal resistant transparent glass. The second

parameter targeted to influence gas absorption was the gas

permeability from the sand mould out to the surroundings.

The front side of the mould was in every case impermeable

due to usage of the glass. To provoke a low permeability, the

back side of the lower mould was covered by a steel plate.

The variation in the permeability was obtained by perforat-

ing the steel plate with an open area of 40%. The mould ma-

terial was made in sand using two different binders, Epoxy-

SO2 and PUCB (Phenolic Urethane Coldbox). The influence

of coating the inner surface of the mould was investigated

by coating with an aluminium silicate based coating. The

chemical composition of the melt was held in the middle of

the variation range reported in Table 1. The average pouring

temperature was 1445°C (2633°F). The measuring methods

of hydrogen concentrations were identical with those used

for measurement in the previous section. The measurement

method of nitrogen concentration for samples taken before

pouring were identical with the previously described meth-

od. The samples aimed for investigation of the nitrogen con-

tent after mould filling were cut out from the open feeder.

Due to the slow solidification rate, the samples was remelted

and cast into a die to obtain white solidification structure

necessary for the optical emission spectroscopy investiga-

tion. The experimental setup is summarized in Table 2.

Figure 1. Mould layout for

measuring gas absorption

Figure 2. Mould layout

surface turbulence.


Measurement of Hydrogen Concentration

Independent of the binder system and other variables, the hy-

drogen concentration in the melt increased by 60-80% during

mould filling. The highest amount of hydrogen measured in the

mould was 3.8 ppm, which is close to the solubility limit of

4 ppm at the solidification temperature of grey iron. If active

nucleation sites are available and the concentration of nitrogen

is high, this level might be dangerous with respect to gas poros-

ity formation. As seen in Fig. 4, the absorption of hydrogen dur-

ing mould filling is affected by the different mould parameters.

When the permeability of the mould increased, the absorption

of hydrogen decreased, which is reasonable as both the air in

the mould cavity and the gases evolved during binder pyrolysis

can escape more easily with increased permeability. It was also

found that the flow velocity, or degree of turbulence, affects the

absorption of hydrogen. The average flow velocity was calcu-

lated from the frames of a high-speed camera observed through

transparent glass. The time necessary for the top front of the

liquid metal to flow from one turning point to the next turning

point of the zigzag trajectory was considered in the calculation

of the average velocity. The lowest change of hydrogen concen-

tration was obtained in the filter mould, which also had the low-

est average filling velocity, 35 cm/s. The highest increase of hy-

drogen concentration was measured in the turbulent mould and

the average filling velocity in this case was approximately 100

cm/s. These values can be compared with the reference mould

where the filling velocity was 55 cm/s. In Fig. 5, the change of

hydrogen is plotted as a function of the average flow velocity

of the molten metal. The natural explanation of these results is

dependent on the filling velocity. The surface turbulence cre-

ates different large interacting surfaces between the flowing

liquid and the gaseous atmosphere. Large velocity associated

to high surface turbulence promotes the hydrogen absorption

mechanism. It can clearly be seen in the high velocity films that

liquid falls back and forms folds in the turbulent mould, while

the flow front is held together in the filter mould. Fig. 3 shows

a still picture from a high velocity film taken during filling of

the turbulent mould. Use of PUCB binder increased the hydro-

gen concentration more than the epoxy binder. In earlier work14

it was shown that PUCB binder evolves more gas than epoxy

binder, at the same corresponding binder contents, and this ex-

plains the results obtained in this case.

Measurement of Nitrogen Concentration

Since the samples from the mould taken for analysis of the

nitrogen concentration needed to be remelted, since they so-

lidified as grey iron and not chilled samples, exact values of

the influence of mould filling could not be obtained. The re-

melting decreased the amount of nitrogen in the samples. The

reference samples taken from the molten iron before mould

filling were not remelted as they had chilled fast and resulted

in white solidification. It seems that the nitrogen concentra-

tion decreases during mould filling, but since this is unlikely,

it is most likely a result of the remelting process. To correct

for this, the results were recalculated by adding the lowest cal-

culated value of the nitrogen change, i.e. 13 ppm to all results,

so the lowest value will be zero and all other results positive.

The recalculated values are seen in Fig. 6. The original ni-

trogen concentration in the iron before mould filling varied

between 27 ppm and 49 ppm. As shown in Fig. 6 the highest

absorption of nitrogen during filling is measured in the mould

with PUCB binder. Since this binder contains nitrogen, in

contrast to epoxy, this is the expected result. As for hydrogen,

there seems to be a relationship between nitrogen pick-up and

filling velocity. This is unexpected because nitrogen porosity

is generally related to gas evolution from nitrogen-containing

binder systems and not absorption of nitrogen from air in the

mould cavity. Decreased permeability seemed to increase the

absorption of nitrogen, which also indicates that nitrogen ac-

tually is absorbed from the air.


Shrinkage Porosity

Shrinkage Porosity in Gray Iron Cast Components

Shrinkage porosity in complex shaped cast iron components,

particularly in cylinder heads, are defects that compromise

the functionality of an engine during the operation. The main

cause for reduced functionality is due to leakage between the

different sections of the cylinder head designed for trans-

porting fuel, cooling media, intake or exhausting gases. A

standard measure to avoid leakage is to apply a Tellurium

based coating on surfaces where leakage is observed. Tel-

lurium, a strong carbide promoter, contributes to the creation

of a thin carbide layer on the casting surface and impedes

the leakage through this surface. The drawback of this appli-

cation is the brittle properties of the carbide with increased

risk for crack initiation. A less active element, Bismuth, is

another element added to coatings to prevent leakage. The

reason for the positive effect of Bismuth is not well under-

stood. The possible influence of Bismuth on the austenite

dendrite growth could affect the dendrite shape.15 From this

observation, the authors conclude that Bismuth influences

the path length between the austenite dendrite arms.

Cylinder heads from two different foundries has been ex-

amined. Both castings, with and without leakage defects,

were included. The scope of the examination was to identify

to all results.

possible factors influencing the shrinkage porosity forma-

tion mechanisms. Both foundries use induction furnaces for

melting the iron with compositions included in the range

presented in Table 1. The castings are molded in green sand

at both foundries, but the internal cores were processed us-


ing different core bounding systems. A basic observation is

that on the surface of the casting identified as an area for

leakage is the ragged surface, in comparison to surfaces

where no leakage was observed. On the surfaces where no

leakage is present, the surface is smooth. Ragged surfaces

include craters (Fig. 7.) with a configuration similar to a na-

ked primary austenite dendrite. The cross section of the cast-

ing in the area with leakage is presented in a color etched mi-

crograph on Fig. 8. In many cases the porosities are clearly

seen to be connected to the casting surface. Fig. 9 indicates

an internal porosity where the cavity is bordered between

primary austenite dendrites and the surface of the dendrite is

covered with thin oxide layers. The oxide layers are believed

to form at elevated temperatures due to contact with oxy-

gen from the atmosphere, outside of the casting. The color

etched picture indicates the distribution of the porosities in

the microstructure exclusively in the neighborhood of the

eutectic cells. In single cases a eutectic cell is surrounded by

porosity, but most often several eutectic cells are engulfed

by porosities. The observations presented clearly note that

the observed leakage was caused by porosities enmeshing

the metallic matrix but it gives no indication on when the po-

rosities are formed during the solidification interval. General

indications in the literature for shrinkage porosity formed at

the end of solidification cannot be confirmed from this ob-

servation due to the larger number of eutectic cells engulfed

by a coherent porosity area.

Studying Shrinkage Porosity Formation

Earlier investigation by the authors16 indicates that macro-

structure of gray cast iron, including austenite grains, contains

several eutectic cells between the primary austenite dendrites.

This observation, in context with coherent shrinkage porosity

engulfing several eutectic cells, may raise the question wheth-

er this type of shrinkage porosity is not formed earlier during

the solidification, being influenced by the macrostructure and

the primary austenite dendrite network.

The complex shaped cast component is not satisfactory to

perform structured investigation including macrostructure

investigation. A heat treatment called DAAS17 is neces-

sary to retain the grain structure of cast iron, otherwise the

austenite will transform to pearlite. For this purpose a spe-

cial sample with a less complex shape than a cylinder head

but including combinations of cylindrical and plate shaped

geometrical components, was designed. (Fig. 10.). The test

samples were molded in furan bounded sand and cast from

the same alloy as was used for the cylinder head investiga-

tion. The samples were DAAS treated and the cross sec-

tion in the bottom of the middle cylindrical bar has been

identified as containing shrinkage porosities of the same

type and at the same extent as was found in cylinder heads.

Fig. 11 shows the grain boundaries of the primary austenite

dendrites. There is a clear distinction between the shape

and the perpendicular position of the columnar austenite

dendrite in relation to the mould wall. There is a layer of

primary columnar dendrites on each surface and between


these layers there is an area containing the equiaxed aus-

tenite dendrites. The shrinkage porosities are found in the

lower right side of the cross section of the sample. At a

higher magnification (Fig 12.) it is clearly visible the posi-

tion of the porosities concentrated in the border between

the columnar austenite grains. The thermal behavior of this

sample is comparable with the thermal distribution of a

cylinder head. An important observation, after solidifica-

tion simulation of the sample, is that the local hot spot of

the sample in the cross section of the middle cylindrical

bar is migrating during solidification. The position of the

hot spot as a function of time and temperature is indicated

in Fig. 13. The simulation software used to perform the so-

lidification simulation was MAGMAsoft. Another impor-

tant observation is the coincidence between the moving hot

spot and the shape and extent of the austenite grains in the

layer to right, which incorporates even the shrinkage po-

rosities. The austenite grains look out to be smaller than on

the opposite wall, but the transition between the columnar

to equiaxed grains is situated at the same distance from the

respective surfaces. A possible explanation is the very slow

and distorted grain growth direction due to the moving hot

spot where the grain boundaries have been changed similar

to a recrystallization process.

The conclusion from the present investigation is that the

time for this type of shrinkage porosity formation is early

in the solidification process, and probably coincides dur-

ing the columnar to equiaxed transition when the need

for feeding the austenite growth leads to transport phe-

nomena between the primary austenite grains. Gases from

the surrounding atmosphere are sucked into the border

grains.

function of space and time.

treatment. Columnar austenite grains


between the columnar austenite dendrite. The distorted

austenite grains, due to the hot spot migration, are be-

lieved to be extremely permeable thus facilitating the

transport phenomena. Once the porosities formed in the

border of the austenite grains, the solidification continues

with the coarsening of the austenite dendrite followed by

the growth of the eutectic phase. In this case, the eutec-

tic phase is developing after the porosities were formed

and the surplus of liquid due to graphite expansion is not

enough to compensate for the existing pores. The pro-

posed mechanism for porosity formation is useful to ex-

plain how the tellurium-based coating was successful in

preventing shrinkage porosity formation.

Gray Cast Iron Cast Components

Collaborating research partner foundries have reported

that metal expansion penetration appears on surfaces

where leakage due to shrinkage porosity occurs. In cyl-

inder heads the most affected surfaces are between the

casting and complex shaped cores forming an internal

cavity for transport of the cooling media (Fig. 14). These

surfaces are often not accessible and only be examined

with difficulty. Non-cleaned surfaces, from the mixture

of a metallic-sand grain matrix, may reduce the cooling

effectiveness and disturb the optimal engine operation.

Examination of cylinder heads and the metallic–sand

grain matrix reveal that the metallic fraction of the pen-

etrated mixture has different microstructure. In minor

cases the matrix includes a mixture of metallic phases

and graphite where the metallic phase contains both pri-

mary and eutectic austenite, which corresponds with the

hypo-eutectic composition used for produce this cast-

ings (Fig. 15). The dominant case is when the metallic

matrix contains pure eutectic phase (Fig. 16), which is

deduced from the appearance of the colour etched mi-

crographs completely missing the primary phase. The

observations indicate that the penetration moment could

happen at different stages of the solidification, such as

during the primary solidification indicated by the hypo-

eutectic composition of the penetrated metal and during

the eutectic solidification indicated by the pure eutectic

composition in the penetrated matrix.

austenite.


For the same reason, as indicated in the case of shrink-

age porosity, to improve the ease of examination, a spe-

cial designed casting sample has been created to investi-

gate the mechanism of expansion penetration formation.

The simulation indicated similarities to the shrinkage

porosity casting samples in that the hot spot of the cast-

ing is migrating from an initial position corresponding

to the mass center of the casting to the surface between

the concave metal surface and the mould wall (Fig. 17).

The test mould has been formed in bonded sand using

Epoxy-SO2 method and DAAS treated. The coincidence

is obvious with the case of samples used to study shrink-

age porosity; the columnar austenite dendrite zone on

the concave casting surface is shorter than in connec-

tion to the other surfaces (Fig. 18). In this case we can

conclude that the migrating hot spot has an influence on

the development of the columnar austenite dendrite in

the area where metal expansion penetration is expected.

Two mainly different mechanisms have been observed

from these experiments. One mechanism is the complete

distortion of the columnar zone (Fig. 19). The area be-

hind the distorted dendrite zone contains small primary

austenite dendrites surrounded by a metallic phase con-

taining carbides (Fig. 20). The conclusion from this ob-

servation is that an internal pressure of the liquid has de-

formed the columnar zone. The concave surface swelled

backward and the area behind was filled up with a liquid

of hypoeutectic composition. The time for this type of

penetration is believed to coincide with the primary so-

lidification. The second observed mechanism indicates

a metal strip, squeezed behind the sand grains forming

the mould surface (Fig. 21). The observed color etched

micrograph indicates this strip being of a pure eutectic

composition without any primary phase (Fig. 22). The

time for this type of penetration is suggested to be dur-

ing the eutectic solidification in the bulk metal when the

surplus of liquid from the interspaces of austenite den-

drite is squeezed between the sand grains.

radius is 15 mm.


Conclusion

A strong comprehensive collaboration between Swedish

foundries including new investigation methods and novel

experimental setup has pointed out some basic understand-

ing on the formation mechanisms of gas porosity, shrink-

age porosity and metal expansion penetration in gray cast

iron components.

New investigation methods used to obtain the results report-

ed in this paper are:

in liquid iron

including the primary austenite phase

-

crostructure on the same micrograph

Novel experimental setups used to obtain the results report-

ed in this paper:

surface for observing the liquid flow

during mould filling

-

rosity and metal expansion penetration

The variety of the tools including methods and experimental

setups, necessary to obtain the presented complex mecha-

nisms are considered to make the difference between witch-

craft and multidisciplinary science.

REFERENCES

1. Dugic, I., “The Mechanism of Metal Expansion

Penetration During Solidification of Gray Cast Iron,”

Linköping Studies in Science and Technology,

Dissertation No. 1007, (2006).

2. Orlenius, J., “Factors Related to the Formation of Gas

Porosity in Gray Cast Iron,” Research Series from

Chalmers University of Technology, ISSN 1653-8891,

Licentiate Theses. (2008).

3. Elmquist, L., “Defect Formation during Solidification

in Gray Cast Iron,” Research Series from Chalmers

University of Technology, PhD Theses. (2009).

4. Dawson, J.V., Kilshaw, J.A., Morgan, A.D., “The

Nature and Origin of Gas Holes in Iron Castings,” AFS

Transactions, vol 73, pp 224-240 (1965).

5. Hernandez, B., Wallace, J.F., “Mechanisms of Pinhole

Formation in Gray Iron,” AFS Transactions, vol 87, pp

335-348 (1979).

6. Dawson, J.V., “Pinhole Defects,” BCIRA, Report 653,

pp 433-437 (1962).

7. Svoboda, J.M., “Fundamentals of Physical Chemistry

Relating to Gases in Cast Metals,” Gases in Cast

Metals; Williamsburg; VA; 24-26 Mar. 1976, 46 pp

(1976).


8. Svensson, I., “Influence of Deoxidation Practice on

Hydrogen Pick-up in Steel Melts,” Solidification

Technology in the Foundry and Cast House; Coventry;

England; 15-17 Sept. 1980, pp 253-262 (1983).

9. Svensson, I., Fredriksson, H.S., “Effect of Hydrogen and

Nitrogen on Formation of Pores in Iron,” Solidification

Technology in the Foundry and Cast House; Coventry;

England; 15-17 Sept. 1980, pp 376-380 (1983).

10. Fruehan, R.J., “Gases in Metals,” ASM Handbook, vol

15, pp 82-87 (1988).

11. Shusen, W., Libiao, W., “The Variation of Hydrogen,

Oxygen, and Nitrogen Levels in Cast Iron with

Varying Blast-Air Moisture,” Cast Metals, vol 5, no 2,

pp 87-91 (1992).

12. Carter, S.F., Evans, W.J., Harkness, J.C., Wallace, J.F.,

“Factors Influencing the Formation of Pinholes in Gray

and Ductile Iron,” AFS Transactions, vol 88, pp 245-

268 (1980).

13. Caspers, K.-H., “Influence of Steel Scrap and Burden

Additives in Induction Furnaces on the Tendency for

Defect Formation in Cast Iron,” AFS International Cast

Metals Journal, vol 2, No 29, pp 29-32 (1977).

14. Orlenius, J., Gotthardsson, U., and Diószeg, A.,

Submitted to International Journal of Cast Metals

Research. (2008).

15. Winardi, L., Loper, C.R., “Influence of Bismuth

Addition on Primary Austenite Dendrite in Gray Cast

Iron”, AFS Transactions, vol 112, pp 723-742 (2004).

16. Diószegi, A., Liu, K.Z., and Svensson, I.L.,

“Inoculation of Primary Austenite in Grey Cast Iron,”

International Journal of Cast Metals Research. Vol. 20,

no. 2, pp. 68-72 (2007).

17. Rivera, J.L., Boeri, R.E., Sikora, J.A., “Solidification

of Gray Cast Iron,” Scripta Materialia, vol.50, pp.331-

335, (2004).

DEFECT FORMATION OF GRAY IRON CASTING Abstract

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