1 - CONTENTS 1 Effects of Alloying Elements on the Microstructures and

Mechanical Properties of Heavy Section Ductile Cast Iron 2 Austempered Ductile Iron Castings for Chassis Applications 3 Austempered ductile iron (ADI) is stronger per unit weight than

aluminium, highly wear resistant and easier to machine than free machining steel

4 Austemperıng Ductıle Iron 5 Fatigue Performance Comparison and Life Predictions of Forged

Steel and Ductile Cast Iron Crankshafts 6 Ductile Iron Microstructure by Thermal Analysis 7 To Study the Effect of Austempering Temperature on Fracture

Behaviour of Ni-Mo Austempered Ductile Iron (ADI) 8 Developments in Cast Iron Metallurgical Treatments 9 GGG 40.3 10 A Review of Common Metallurgical Defects in Ductile Cast Iron 11 Nodular Cast Iron Fatıgue Lıfetıme In Ultra-Hıgh-Cycle Regıon 12 Nodulizing and Inoculation Approaches for Year 2000 and

Beyond - Part 1 13 Some Studies of Nodular Graphite Cast Iron 14 Suggestions for Improved Reliability in Thermal Analysis of Cast

Irons

J. Mater. Sci. Technol., Vol.23 No.l, 2007 97

Effects of Alloying Elements on the Microstructures and Mechanical Properties of Heavy Section Ductile Cast Iron

G.S.Cho1^, K.H.Choe^, K W.Lee 1) and A.Ikenaga2^ 1) Advanced Material Processing Team, Korea Institute of Industrial Technology, 994-32, Dongchun-Dong, Yeonsu-Ku,

Incheon 406-130, South Korea 2) Department of Metallurgy and Materials Science, Graduate School of Engineering, Osaka Prefecture University, 1-1,

Gakuen-Cho, Sakai, Osaka 599-8531, Japan [Manuscript received February 17, 2006, in revised form May 29, 2006]

The effects of alloying elements on the as-cast microstructures and mechanical properties of heavy section ductile cast iron were investigated to develop press die material having high strength and high ductility. Measurements of ultimate tensile strength, 0.2% proof strength, elongation and unnotched Charpy impact energy are presented as a function of alloy amounts within 0.25 to 0.75 wt pet range. Hardness is measured on the broken tensile specimens. The small additions of Mo, Cu, Ni and Cr changed the as-cast mechanical properties owing to the different as-cast matrix microstructures. The ferrite matrix of Mo and Ni alloyed cast iron exhibits low strength and hardness as well as high elongation and impact energy. The increase in Mo and Ni contents developed some fractions of pearlite structures near the austenite eutectic cell boundaries, which caused the elongation and impact energy to drop in a small range. Adding Cu and Cr elements rapidly changed the ferrite matrix into pearlite matrix, so strength and hardness were significantly increased. As more Mo and Cr were added, the size and fraction of primary carbides in the eutectic cell boundaries increased through the segregation of these elements into the intercellular boundaries.

K E Y W O R D S : Heavy section ductile cast iron; Alloying elements; As-cast microstructures; As-cast mechanical properties

1. Introduction

Ductile cast iron exhibits good ductility and toughness because the graphite morphology is spher­ical. It has been reported that the strengthening and toughening of ductile cast iron result from the modification of the matrix structure when alloying elements'1] are added or heat treatment^ is applied. The austempered ductile cast irons have been studied to replace the forged or cast steel in the structural parts, because they have many advantages such as high strength and toughness, high machinability and good corrosion resistance I 3 - 6 ! . Recently, the heavy section drawing dies in pressing steel sheet for the automobile frame or panel are replaced by simple one body-type as-cast ductile cast iron, which yields low cost and shortened delivering time on producing press dies. The casting die is generally produced via a full mold process that uses the near-net shaped ex­panded polystyrene pattern replaced by the molten m e t a l A s - c a s t ductile cast iron should meet the re­quirements of material properties to be used in cold pressing dies including high strength and high ductil­ity, but both mechanical properties are very difficult to obtain within the same casting material. Thus, the amounts of alloying elements are controlled to achieve as-cast heavy section ductile cast having high strength and ductility. The purpose of this paper is to obtain experimental data for optimum alloy design of heavy section ductile cast irons and to present the effects of alloying elements on the as-cast microstructures and mechanical properties for using cold press die mate­rial.

f Ph.D., to whom correspondence should be addressed, E-mail: gscho@kitech.re.kr.

2. Experimental

The chemical compositions of ductile cast irons used for this work are presented in Table 1. The main difference lies on the Mo, Cu, Ni and Cr al­loy contents. The nominal compositions of alloy­ing element additions are given as 0.25, 0.5 and 0.75 wt pet. The results are presented in terms of the nominal additions for convenience in discus­sion. The ductile irons were obtained by melting steel scrap, pig iron, graphite, Fe-75 wt pet Mn and Fe-75 wt pet Si in the high frequency induction fur­nace. Spheroidizing and inoculation practices were performed in a conventional sandwich method with 5.8 wt pet Mg-Fe-Si and 75 wt pet Si-Fe alloy, respec­tively. The metal was poured into furan resin molding sand molds to obtain Y-shaped 75 mm blocks. Ten­sile specimens with the dimensions shown in Fig.l and

200

75

<—•

Fig . l Y-shaped block and tensile specimen dimensions

98 J. Mater. Sci. Technol., Vol.23 No.l, 2007

Table 1 Chemical compositions of ductile cast iron (wt pet)

Melt C Si Mn P S Mo Cu Ni Cr Mg A 3.61 2.36 0.28 0.04 0.01 - - - - 0.06 B 3.57 2.37 0.28 0.04 0.01 0.19/0.38/0.58 - - - 0.06 C 3.60 2.40 0.28 0.04 0.01 - 0.24/0.47/0.70 - - 0.06 D 3.64 2.35 0.30 0.04 0.01 - - 0.24/0.48/0.70 - 0.06 E 3.66 2.31 0.33 0.04 0.01 - - - 0.26/0.49 0.06

Table 2 Typical characteristics of spherical graphite in ductile cast iron

Melt Area fraction/% Nodule count/(N-mm 2) Nodularity/(%) Ave. diameter/jum A 10.2 141.0 90.4 30.3 B 13.1 153.3 88.5 33.0 C 13.0 154.6 87.5 32.8 D 12.2 152.6 87.3 31.8 E 11.4 148.0 86.0 31.5

Fig. 2 As-cast microstructure of melt A as a reference casting

unnotched Charpy specimens (55 mmx 10 mmx 10 mm) were machined from the bottom section of the Y-blocks. Tensile specimens were taken from the bottom section of Y-blocks in order to minimize casting de­fects such as microporosity. Three tensile specimens were tested in a 250 kN hydraulic Instron universal testing machine using a constant cross-head travel speed of 4 mm/min. Unnotched Charpy specimens were tested in a standard impact testing machine. Hardness measurements were made on a standard Brinell hardness machine with a load of 3000 kg and 10 mm steel ball. Metallographic specimens were obtained from broken impact samples.

3. Results and Discussion

3.1 As-cast microstructures of ductile cast iron Figure 2 shows the as-cast microstructure of melt

A as a reference cast iron given in Table 1. The mi­crographs show that the matrix structure is fully fer¬rite structure with a small fraction of pearlite near the eutectic cell boundaries. The ferrite structures are selected to serve as a base material by minimiz­ing Mn content in 75 mm heavy section Y-block. The spheroidal graphite nodules are well distributed within the ferrite matrix. Two types of graphite in size are observed microscopically, because of the slight hypereutectic chemical composition of the reference cast iron. The bigger graphites, about 50 /xm in diam­eter, are crystallized at the graphite+liquid region in the Fe-C system. The smaller ones, under 50 /xm, are

formed through eutectic solidification range. Image analysis of volume fraction, nodule count, nodular­ity and average diameter of graphite on the different alloyed ductile cast irons are given in Table 2. The characteristic values of graphites in the alloyed irons are very similar to that of reference cast iron.

Figure 3 shows the as-cast microstructures of Mo alloyed cast iron. The Mo alloyed ductile cast iron of melt B exhibits substantial ferrite in the as-cast structure with some pearlite at intercellular regions. As more Mo is added, the matrix becomes fully fer­rite and the area fraction of pearlite structure in the eutectic cell boundaries is slightly increased. The Ni alloyed ductile cast iron of melt D also has a mostly ferrite matrix having some pearlite structures in the cell boundaries. Adding up to 0.75 wt pet of Mo and Ni elements results in the same tendency on the for­mation of a ferrite matrix. The addition of Mo and Ni alloy will increase the hardenability of ductile cast iron by delaying the transformation of austenite to ferrite^8,9!. Also, the hardness of the ferrite matrix increased from about 169 HV to about 188 HV when more Mo was added. The Ni alloyed ductile cast iron also exhibited an increase in matrix hardness. This means that the addition of Mo and Ni strengthens the ferrite matrix via solid solution hardening.

Figure 4 shows microstructures of the as-cast melt C ductile cast iron. As more Cu is added, the amount of ferrite structure is significantly decreased with Cu content. For the melt C containing 0.5 wt pet Cu and more, the matrix was abruptly changed into pearlite

J. Mater. Sci. Technol., Vol.23 No.l, 2007 99

Fig.3 Microstructures of the as-cast melt B ductile cast iron with different Mo contents: (a) 0.25 wt pet Mo, (b) 0.5 wt pet Mo, (c) 0.75 wt pet Mo

Fig.4 Microstructures of the as-cast melt C ductile cast iron with different Cu contents: (a) 0.25 wt pet Cu, (b) 0.5 wt pet Cu, (c) 0.75 wt pet Cu

IMB

Fig.5 Microstructures of the as-cast ductile cast irons with different alloying elements: (a) 0.75 wt pet Mo (b) 0.75 wt pet Cu, (c) 0.75 wt pet Ni, (d) 0.5 wt pet Cr

with a bull's-eye ferrite rim around the nodules. Fig­ure 4 presents the ability of Cu to hinder the diffu­sion of carbon into graphite at the graphite-austenite interface during the solid state transformation after solidification. The Cu will increase as-cast strength and hardness through increased pearlite formation!8].

Figure 5 shows the as-cast microstructures of duc­tile cast iron with different alloying elements. While

the 0.75 wt pet Mo and 0.75 wt pet Ni alloyed ductile cast irons exhibit a full ferrite matrix, the 0.75 wt pet Cu and 0.5 wt pet Cr added ductile cast irons ex­hibit a bull's-eye structure within the pearlite matrix. Carbide-like phases are observed in the middle of the eutectic cell boundaries of the 0.75 wt pet Mo and 0.5 wt pet Cr alloyed ductile cast iron. These phases are not observed at the intercellular regions of

100 J. Mater. Sei. Techno!., Vol.23 No.l, 2007

Mn « Fig.6 EDX line scan profiles measured on as-cast ductile cast irons: (a) 0.75 wt pet Mo, (b) 0.5 wt pet Cr

Alloy content / wt pet Alloy content / wt pet

Fig.7 Variation in tensile strength and elongation of as-cast ductile cast irons having different alloying contents: (a) tensile strength, (b) elongation

0.25 wt pet Mo and 0.25 wt pet Cr alloyed ductile cast iron. Figure 6 presents the E D X line scan profiles measured on the as-cast ductile cast irons. Figure 6(a) shows the sharp increase in Mo peak at the eutectic cell boundary. This means that the segregation of Mo during solidification forms Mo-rich carbides in the cell boundaries^10"12]. Figure 6(b) shows the segregated Cr-rich phase in the as-cast ductile cast iron. Mn was positively segregated in the cell boundaries but Si was negatively segregated^13!. The increased Mo and Mn levels in the intercellular boundary in the alloyed iron attributed to segregation caused considerable pearlite formation near the cell boundaries. Carbide forma­tion will deteriorate the ductility of as-cast ductile cast front1 4'1 5!. Thus, the amount of Mo and Cr addi­tions should be limited to a certain level considering the thickness of castings.

3.2 Mechanical properties of as-cast ductile cast iron Figure 7 shows the variation in tensile strength

and elongation of as-cast ductile cast iron having dif­ferent alloying contents. The trends in terms of tensile properties of Ni added ductile cast iron resemble that of Mo added ductile iron. With the increase in the amount of Mo and Ni additions, tensile strength is slightly increased while elongation is decreased. The strength of Mo and Ni alloyed ductile cast iron is mainly dependent on the solid solution hardening of ferrite matrix. It was also considered that the de­crease in elongation was caused by the increase in pearlite formation in the cell boundaries. As more Cu element is added, tensile strength is steeply increased. The tensile strength of 0.5 wt pet Cu added ductile cast iron approaches at a maximum of 700 MPa, but elongation quickly decreases below 5%. The trends in terms of tensile properties of Cr added ductile iron

resemble that of Cu added ductile cast iron. Addi­tion of up to 0.5 wt pet Cr could not reach a maxi­mum of 700 MPa because of the formation of a large bull's-eye ferrite structure. It is confirmed that the tensile strength of as-cast ductile cast iron is strongly dependent on the area fraction of pearlite, while elon­gation is inversely proportional to the pearlite frac­tion in the as-cast microstructuret 1 6 - 1 8]. Figure 8 presents the S E M micrographs of the fractured sur­face of tensile specimens for comparing two typical fracture modes. Figure 8(a) shows the brittle fracture surface of 0.75 wt pet Cu alloyed iron having pearlite matrix with cleavage river patterns. The ductile frac­ture surface of 0.75 wt pet Ni alloyed iron is shown in Fig.8(b). Plastic deformation of ferrite around the boundary of graphite nodule is observed. Some cleav­age brittle fracture area having a pearlite phase near the intercellular region is also shown.

Figure 9 shows the variation in hardness and im­pact energy with different alloying additions. As more alloying contents are added, the Brinell hardness of Mo and Ni alloyed ductile cast iron is linearly in­creased, but that of Cu and Cr alloyed ductile cast iron increased very steeply. The Brinell hardness of ductile cast iron refers to the resistance of the matrix and graphite to plastic deformation. The variation in hardness with the alloying element is strongly de­termined by the as-cast microstructures. Generally, the hardness of cast iron is increased by the volume fraction of pearlite structure^19'20!. In this study, the addition of Mo and Ni slightly increased the matrix hardness via solid solution hardening. As more Cu and Cr element is added, the hardness increased sig­nificantly owing to the high fraction of pearlite. The Charpy impact energy of Mo and Ni alloyed iron is gradually decreased, but that of the Cu and Cr

Fig.8 SEM micrographs showing the fractured surface of tensile specimens: (a) 0.75 wt pet Cu, (b) 0.75 wt pet Ni

1 0 0 | 1 , 1 , 1 , 1 1 0 I 1 • 1 1 1 • 1 1

0 0.25 0.5 0.75% ° 0.5 0.75

Alloy content / wt pet Alloy content / wt pet

Fig.9 Variation in hardness and impact energy of as-cast ductile cast irons having different alloying contents: (a) Brinell hardness, (b) Charpy impact energy

alloyed iron begins to decrease fast. This indicates that ferrite matrix provides highly ductile cast iron while pearlite yields highly strong cast iron. Mo was found to reportedly segregate at cell boundaries and aggravate the toughness of ductile cast iron! 2 1 !. How­ever, the toughness of ductile cast iron was more de­pendent on the area fraction of ferrite than the area fraction of intercellular carbide-like phases.

4. Conclusions

(1) The as-cast microstructures are strongly de­pendent on the alloying elements. The ferrite matrix is mostly formed as Mo and Ni elements are added. The addition of Cu and Cr rapidly changed the ferrite into pearlite matrix.

(2) The Mo and Cr element are segregated at the eutectic cell boundaries and caused the formation of carbide-like phases. Therefore, the amount of Mo and Cr added should be controlled considering the thick­ness of castings.

(3) The ferrite matrix of Mo and Ni alloyed duc­tile cast iron exhibited low strength and low hardness but high elongation and high impact energy. The me­chanical properties of pearlite cast iron alloyed wi th Cu and Cr element were the exact opposite of those of ferrite cast iron.

(4) Ferrite matrix provides highly ductile cast iron while pearlite provides highly strong cast iron. The toughness of ductile cast iron was more dependent on the area fraction of ferrite than the area fraction of intercellular carbide-like phases.

REFERENCES

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[2 ] R.C.Voigt and C.R.Loper: J. Heat Treating, 1984, 3, 26.

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[4 ] B.V.Kovacs: Mod. Cas., 1987, 35, 314. [5 j C.S.Choy, W.Sharpe, J.Barker and F.J.Fields: Metall.

Mater. Trans., 1996, 27A, 923. [6 ] O.Eric, L.Sidjanin, Z.Miskovic, S.Zec and

M.T.Jovanovic: Mater. Lett., 2004, 58, 2707. [7 ] Susan P.Thomas: Expendable Pattern Casting, Vol.1,

American Foundrymen's Society, Inc., 1994. [8 ] M.F.Burditt: Ductile Iron Handbook, AFS, Des

Plaines, IL, 1993, 63. [9 ] M.N.Ahmadabadi: Metall. Mater. Trans., 1997, 28A,

2159. [10] H.Bayati and R.Elliott: Mater. Sci. Technol, 1995,

11, 118. [11] J.Liu and R.Elliott: Int. J. Cast Metais Res., 1999,

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Met, 1994, 4, 182. [13] A.Owhadi, J.Hedjazi, P.Davami, M.Fazli and

J.M.Shabestari: Mater. Sci. Technol, 1997, 13, 813.

[14] E.Dorazil: High Strength Austempered Ductile Cast Iron, 1 s t ed. London, Ellis Horwood, 1991.

[15] M.Caldera, G.L.Rivera, R.E.Boeri and J.A.Sikora: Mater. Sci. Technol, 2005, 21(10), 1187.

[16] E.Carren, M.Diao and R.Schaller: Sci. Mater., 1997, 38, 259.

[17] G.Wilkinson and C.Grupke: J. Heat Treating, Second Int. Conf. on Ductile Iron, Ann Arbor, MI , 1986, 349.

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[19] C.Ji and S.Zhu: Mater. Sci. Eng., 2006, 8A419, 318. [20] Charls F.Walton: Iron Castings Handbook, Iron Cast­

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Trans., 1984, 92, 801.

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SAE TECHNICALPAPER SERIES 2000-01-1290

Austempered Ductile Iron Castings forChassis Applications

Robert J. Warrick, Paul Althoff, Alan P. Druschitz,Jeffrey P. Lemke and Kevin Zimmerman

Intermet Corp.

P. H. ManiDuctile Iron Society

Mitchell L. RackersCaterpillar Corp.

Reprinted From: Casting Solutions for the Automotive Industry(SP–1504)

SAE 2000 World CongressDetroit, MichiganMarch 6–9, 2000

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1

2000-01-1290

Austempered Ductile Iron Castings for Chassis Applications

Robert J. Warrick, Paul Althoff, Alan P. Druschitz, Jeffrey P. Lemke and Kevin ZimmermanIntermet Corp.

P. H. ManiDuctile Iron Society

Mitchell L. RackersCaterpillar Corp.

Copyright © 2000 Society of Automotive Engineers, Inc.

ABSTRACT

Austempered ductile iron (ADI) castings provide a uniquecombination of high strength and toughness coupled withexcellent design flexibility for chassis applications.

This paper describes the development of the uppercontrol arm for the 1999 Ford Mustang Cobra as anaustempered ductile iron casting. The full servicedevelopment process used is described from initialdesign through finite element analysis (FEA), designverification, casting production, heat treatment,nondestructive evaluation and machining. To achievesignificant weight savings, an austempered ductile ironcasting was chosen for this application instead of an as-cast SAE J4341, Grade D4512 ductile iron casting or asteel forging.

This is believed to be the first application of anaustempered ductile iron casting for a safety critical,automotive chassis application.

INTRODUCTION

For the 1999 Ford Mustang Cobra, the initial materialcandidates for the upper control arm (UCA) of theindependent rear suspension (IRS) were an aluminumcasting and a steel forging. A major challenge associatedwith the IRS Mustang was packaging the suspensionwithin the existing body. The tight packagingrequirements, which limited the cross section size,coupled with the load carrying requirements resulted inthe elimination of aluminum as a material candidate. Thepackaging requirements also led to the demise of thesteel forging. The available envelope required a relativelyintricate part shape to satisfy the conflicting criteria ofload carrying and weight requirements and clearanceneeds. The shape requirements clearly favored a castingover a forging.

Ford approached the Intermet Wagner Foundryregarding the design and production of an austemperedductile iron (ADI) upper control arm. After reviewing initialpackaging and strength requirements and a number ofADI concerns (primarily cost related), the challenge wasaccepted.

The subsequent development of the 1999 Ford MustangUCA as an austempered ductile iron casting is describedin this paper. The full service development process usedis described from initial design through machining. Theselection of an ADI casting, instead of an SAE J4341

D4512 ductile iron casting or a steel forging, was found tobe highly beneficial by permitting significant componentweight savings.

The use of ADI castings in vehicle applications has beenremarkably slow as a result of reports of processinconsistency, from both a foundry and heat treatingstandpoint; machinability concerns; and cost unknowns(largely based on lack of production experience forautomotive type volumes). For this program,responsibility for the ADI part from design throughdelivery of the assembled upper control arm was placedwith a single supplier, greatly simplifying the developmentprocess. The results of this program demonstrate thataustempered ductile iron castings can be successfullyused for safety critical, automotive chassis applications.

SELECTION OF ADI FOR CHASSIS APPLICATIONS

Prime requirements for the upper control arm for theMustang Cobra rear suspension were light weightcoupled with strength and toughness.

One obvious reason for looking at a ductile ironcomponent versus a steel design is that, because of thevolume of graphite nodules present, a ductile iron castingwill weigh approximately 10% less than a steel forging ifboth have exactly the same shape. In addition, the ductile

2

iron casting will provide greater shape flexibility, as theresult of reduced draft angle requirements and the abilityto cast in features through the use of cores.

The geometric flexibility of a casting is too oftenunderstated. Beyond satisfying difficult clearance needs,one is often able to incorporate a number of intricatestructural design elements (ribs, webbing, and windows)to maximize the amount of material in structural regionsand reduce the amount of material in less structuralregions.

Figure 1. Minimum properties of conventional and austempered ductile irons as a function of the average hardness for each grade.

The density and design flexibility advantages would notbe sufficient, particularly for light weight designs, if thematerial strength and toughness were inadequate. This iswhere the unique physical properties of austemperedductile iron come into the picture. While conventionalductile iron grades have long been used successfully forautomotive safety components such as steeringknuckles, lower control arms, brake calipers and anchors,the high strength and toughness of austempered ductileiron offers even greater potential for weight reduceddesigns.

The relationships between hardness and yield strength,tensile strength, and elongation for conventional andaustempered ductile are illustrated in Figure 1. In thisfigure, specification1,2,3 minimum values of yieldstrength, tensile strength, and elongation are plottedagainst typical average grade hardness. As expected, forboth conventional and austempered ductile irons, yieldand tensile strength increase and elongation decreases,with increasing hardness. Of particular interest are twokey factors for austempered ductile iron:

1. The magnitude of the yield strength of ADI startsapproximately where the strength of the conventionalductile iron grades leaves off.

2. The elongation of the softest ADI grade is muchhigher than that of the harder and strongerconventional ductile iron grades.

The softest of the austempered ductile iron grades is theone normally considered for chassis applications as aresult of its excellent strength and toughness. In additionto yield and tensile strength, the fatigue strength of ADI ismuch higher than that of the conventional ductile irongrades normally used for chassis applications. This isillustrated in Table 1. The endurance limit of a Grade 1ADI is approximately twice as great as that of aconventional D4512 grade of ductile iron. Further, thefatigue strength of ADI can be greatly increased bymechanical working of the surface layers of the materialthrough such processes as fillet rolling or shot peening.The major strengthening effect of fillet rolling is illustratedthrough the data in Table 1. The minimum specifiedunnotched Charpy impact strength for an ASTM 897Grade 1 ADI is 100 Joules at 22°C ± 4°C (75 foot-poundsforce at 72°F ± 7°F). Austempered ductile irons provide aunique combination of strength and toughness in areadily cast material.

Based on properties, ADI should be widely used and, infact, the rate of use is now increasing quite rapidly;primarily in lower volume non-automotive applications.However, the lack of positive production experience withADI in higher volume automotive applications has limitedthe acceptance of this material by the automotiveindustry. There are three major concerns:

1. Machinability. The lack of volume experience is amajor factor. Best tools and tooling conditions are notwell known. As a result, people have tended to shyaway from the material because of cost andproduction rate unknowns.

Table 1. Approximate Endurance Limit Values for Ductile and Austempered Ductile Irons

MaterialHardness,

BrinellEndurance

Limit, MPa(psi)

Ductile Iron 156 207 (30,000)

Ductile Iron 187 228 (33,000)

Ductile Iron 261 310 (45,000)

Ductile Iron, Fillet Rolled (Production)

187 317 (46,000)

Ductile Iron, Fillet Rolled (Lab)

187 372 (54,000)

ADI 302 414 (60,000)

ADI, Fillet Rolled (Crankshaft)

302 1,000 (145,000)

3

2. Process Inconsistency Reports. Particularly early inADI’s history, there were a number of reports ofmaterial inconsistencies over time; for example, onelot machines fine and the next lot is almostimpossible to machine.

3. Costs. ADI is produced by casting heat treatment, aswill be discussed in more detail later in this report.Also, added alloy is typically required to insure thatundesirable microstructure constituents don’tdevelop during the quench from the austenitizingtemperature to the austempering temperature. Bothfactors increase piece cost relative to conventional,non-heat treated ductile iron castings. However, thelargest cost issue probably relates to uncertaintyabout machining, scrap, and process control costs.

Development work6 has shown that ADI is quitemachinable, can be produced consistently by a good,knowledgeable foundry and heat treater, and that costsare predictable.

THE APPLICATION – For 1999, Ford introduced amodular, independent rear suspension, initially availableonly on the limited-edition Special Vehicle Team (SVT)Cobra, to provide superior performance and handlingrelative to earlier versions. This new independent rearsuspension features a wishbone-type, short and longarm suspension design. It is a self contained, bolt-in rearend that has been discussed in a number of recentarticles4,5. It is built by Benteler Automotive and isillustrated in Figure 2.

Figure 2. Photograph of the independent rear suspension for the 1999 Ford Mustang Cobra.

Adding an IRS would be expected to bring a relativelyhigh weight penalty. This anticipation, along with theprojected increase in vehicle performance, were themajor driving forces for Ford SVT engineers in their questfor tough, high strength, light weight components,including the upper control arms.

The challenge was to design a minimum weight part thatcould be reliably produced under routine productionconditions using a relatively untried material undervolume production conditions. While Intermet hadpreviously produced over 500,000 selectivelyaustempered automotive camshafts and over a millionaustempered compressor crankshafts, this was thecompany’s first production ADI venture for an automotivesafety part.

THE APPROACH – Key advantages from the start of thisprogram were the solid support of Ford SVT engineersand having full responsibility for the upper control armfrom design through delivery of the assembled uppercontrol arm. It’s far easier to control product consistencywhen one has total responsibility for the design and allaspects of production of the component.

A team was brought together to handle various aspectsof the program from component design throughcomponent delivery. All design, casting production, andthe majority of the materials evaluation work were doneinternally at Intermet. Much of the component testingwork was done at Defiance while Ford and Bentelerhandled module and vehicle testing. Applied Process, Incwas selected as the heat treat source for this programsince there was not sufficient internal austemperingcapacity. Climate Control, Inc. an excellent andprogressive machining source near the producingfoundry in Decatur, Illinois was selected as the machiningsource. They worked closely with Wagner Foundrypersonnel to develop the machining strategy for the ADIupper control arm. This close interaction betweendesigners, producing foundry, and machining source washighly beneficial.

It was this team, working together, that made the ADIMustang upper control arm program a success.

COMPONENT DESIGN – The design information initiallyprovided by Ford included wireframe and surfaceinformation for a cast design, which was the final iterationof earlier forged steel designs, and the initial load casesbased on ADAMS results. Starting with this information,the designers went through four design phases.

Figure 3. Two views of the Phase 1 design.

Phase 1 – This design is shown in Figure 3. It isessentially the last forging design modified to be anaustempered ductile iron casting. While approximately10% weight savings can be achieved simply bysubstitution of ductile iron for steel, this phase wasnecessitated by test casting timing requirements ratherthan by weight reduction goals. It also provided a startingpoint for analyses and design improvements. The Phase1 casting met strength objectives but, as expected, was

4

too heavy. The unmachined Phase 1 casting weight was11.43 pounds.

Phase 2 – The Phase 2 design, illustrated in Figure 4, isbasically a weight reduced version of the Phase 1 design.While the casting weight was reduced by 1.4 pounds, thecasting had insufficient strength. It weighed 9.94 pounds.

Phases 1 and 2 were workhorse phases. After these, theloading requirements were revised and additionalclearance issues were identified which further restrictedthe design envelope. The upper control arm now neededto carry additional load with less section.

Figure 4. Two views of the Phase 2 design.

Figure 5. Two views of the Phase 3 design.

Phase 3 – As a result of the increased loadingrequirements and the unsatisfactory overall strength-to-weight relationship for the Phase 2 design, a new shapeapproach was taken for Phase 3, as illustrated in Figure5. Considerable enhancements were made to provide amore rigid structure with better overall stress distribution.(Note the webbing between the forward and aft legs andthe somewhat parabolic shape approaches that of anideal I-Beam or C-Channel with uniform stressdistribution.) This design met the new componentstrength requirements while providing significantly

reduced weight. Preliminary design verification studieswere performed on the Phase 3 design with satisfactoryresults. At this point a final design was close but a fewchanges were deemed desirable to aid manufacturabilityand additional weight reduction opportunities wereidentified. The Phase 3 casting weighed 8.91 pounds.

Phase 4 – During Phase 3 prototype production, adespruing (gating and risering system removal) problemwas identified that frequently caused excessive castingdistortion. For the Phase 4 design, which is shown inFigure 6, a tie bar was added to prevent this problem. Inaddition, clamp pads were added at the ball joint end toprovide greater part rigidity during machining. With thesemodifications for manufacturability, the design was readyfor final verification and subsequent productionverification studies. The Phase 4 casting weight is 8.53pounds.

Initial design work was done using PDGS for the conceptwork, ARIES for pre- and post-processing, andNASTRAN for all finite element analysis work. At present,CATIA, UNIGRAPHICS, and SDRC are used internallyfor most design work along with PATRAN/NASTRAN forFEA analyses and MAGMA for casting processsimulation studies.

Figure 6. Two views of the Phase 4 design.

DESIGN VALIDATION – Design verificationresponsibilities were split between Ford, Benteler, andIntermet. Ford and Benteler conducted all module andvehicle tests. Ford also handled the simulated vehicledurability and load vs. deflection tests. The latter wereconducted to determine any possible interferencebetween the upper control arm assembly and attachingcomponents under maximum road load conditions.Intermet was responsible for the finite element analysisstudies as well as all material tests, salt spray tests,component fatigue, cold impact, and spike stop tests,component tensile and compression tests, and bushingpush-in and push-out tests. A number of the latter testswere performed at Defiance - STS. All design validation

5

tests were passed satisfactorily. Ongoing In-Processtests for component fatigue, component ultimate tension,and bushing push-in and push-out are the responsibilityof Intermet.

The design of the independent rear suspension is suchthat other components will fail before the upper controlarm. The ability of the lightweight ADI upper control armto readily pass 100,000 cycles at 115% of the peak one-event load is a testament to the design potential of ADIfor chassis applications.

CASTING PRODUCTION – Molds for the Ford Mustangupper control arm castings are made on a DISA 2013MK5 B machine. No cores are required. The ductile baseiron for the castings is melted in coreless inductionfurnaces and magnesium treated using the Fischerconverter process. Internally developed inoculationpractices are used along with a pressure pour furnace forfinal metal treatment and pouring.

Keys to successful ADI production include excellentmetal chemistry and inoculation control at the foundry.The amount of carbon in austenite at the time of thequench into molten salt has a significant effect on theproperties developed in the austempered ductile iron.The carbon content of this austenite is controlled by bothmetal chemistry and austenitizing temperature. Since thecastings will be heat treated in lots, consistent chemistrycontrol within each lot is essential for consistent ADIcharacteristics. Also, lot-to-lot chemistry consistency isimportant to minimize the heat treat temperatureadjustments that have to be made between lots. Theeffectiveness and consistency of the internally developedinoculation practice, coupled with the pouringtemperature control resulting from pressure pouring, helpproduce castings with a consistent, relatively high nodulecount. This reduces alloy segregation and thus improvesmicrostructure homogeneity.

As a result of the intentional lack of heavy sections in thefinal upper control arm design, relatively little alloy isrequired to avoid the formation of any pearlite during theaustempering quench operation. While some copper isadded, additions of expensive alloy elements like nickeland/or molybdenum are not necessary.

Inspection operations following casting cleaning includex-ray audits for casting integrity and 100% ultrasonicvelocity inspection for nodularity.

CASTING HEAT TREATMENT – A typical austemperingheat treatment cycle is illustrated schematically in Figure7. The casting is heated to a temperature, typically in the870 to 900° C (1600 to 1650°F) range, and held at leastlong enough to be uniformly at temperature. This can bedone in molten salt or in an atmosphere furnace. Thelatter, with an endothermic gas atmosphere, is used for

the upper control arm. The part is next quenched inmolten salt and held in the salt bath at a temperature,typically in the 260 to 371°C (500 to 700°F) range, for atime generally in the 1.5 to 2 hour range. The actualtemperature is dependent on the casting propertiesdesired. Lower temperatures give higher hardness andstrength. Higher temperatures yield higher toughnessand machinability. For the upper control arm,approximately 377°C (710°F) is used.

Figure 7. Schematic austempering heat treatment cycle.

Since the optimum heat treat temperatures and times arechemistry dependent, relevant casting chemistryinformation is supplied to the heat treater along with thecastings. However, casting chemistries are typically heldtightly enough that heat treat cycle adjustments are notnecessary.

AUSTEMPERED CASTING MATERIALPROPERTIES – The bulk material properties of thecasting are fully established after austempering. Thesurface properties can be further enhanced throughsubsequent operations such as shot peening or filletrolling.

Average properties for test and production lots ofcastings cast over a twelve-month period are shown inTable 2. With excellent toughness, the yield and tensilestrengths of the ADI are at least twice as large as thoseof a conventional SAE D4512 ductile iron. Studies6 showthat the endurance limit of a Grade 1 ADI is alsoapproximately twice that of a conventional D4512 ductileiron.

The ability to achieve excellent strength and toughness,under routine production conditions, makes ADI anexcellent material for chassis components.

6

POST HEAT TREATING/PRE-MACHININGOPERATIONS – Following the austempering heattreatment, the castings are shot peened to removesurface scale and discoloration caused by the heattreatment operation. This operation has the addedbenefit of significantly increasing the fatigue strength ofunmachined casting surfaces. Almen strips were run withthe pre-production approval samples to quantify theamount of peening that the castings were receiving.Almen strips will be run with future shot peening loads asan ongoing quality check.

After peening, the castings are 100% inspected using amultifrequency eddy current unit to verify that eachcasting was properly austempered. This procedure andthe resulting correlations were developed and validatedinternally. At this stage, 100% casting inspection is beingused to help insure that each casting was properlyaustempered and to look for variations that might relateto machining performance. It is important to insure thatall process controls consistently do the necessary job.Also, the additional information being gathered shouldhelp in further optimizing quality and manufacturingperformance.

MACHINING AND ASSEMBLY – Following heattreatment and painting with e-coat, the upper controlarms are machined, gauged, and assembled at ClimateControl. Figure 8 shows two views of a machined uppercontrol arm with the bushings in place. The parts aremachined on heavy duty, horizontal machining centers,as shown in Figure 9. After machining, the parts are100% gauged in-cycle and the bushings are thenpushed-in in a separate station. The 100% gaugingcurrently being done was incorporated into the machiningoperation as part of the ongoing learning process.

Figure 8. Two views of a machined upper control arm casting with the bushings in place.

Figure 9. Photograph of a machining center with the doors open to show the fixture and casting orientation for machining.

The fixtures holding the casting during machining weredeliberately designed to clamp the part solidly to a heavybase to minimize vibration during machining. Aphotograph of a casting in the machining fixture ispresented in Figure 10. Climate Control personnelindicate that a solid setup which minimizes vibrationduring machining is key to successful ADI machining.They further indicate that:

1. They have had no major problems machining the ADIupper control arms.

2. Tool usage is higher than for conventional grades ofductile iron (which are typically much softer).

3. They typically use slower and deeper cuts for ADIthan for conventional ductile irons.

4. They are still evaluating alternate specialty insertgrades to find the best insert for their machiningconditions.

5. A key to successful ADI machining is a stablemanufacturing process for the ADI castings (bothfoundry and heat treater).

Table 2. Comparison of Specification and Average ADI UCA Mechanical Properties

PropertySpecification(ASTM A897/

897M)

AverageADI UCA

Hardness,Brinell

269-321 301

Tensile Strength, MPa (psi)

850 min.(125,000 min.)

1,049(152,100)

Yield Strength,MPa (psi)

550 min.(80,000 min.)

798(115,700)

Elongation,%

10 min. 13.8

Impact Strength, Joules (ft-lbs)

100 min.(75 min.)

179(132)

7

There was an issue early on relating to ability to holddimensions that was traced to casting distortion duringheat treatment. This was resolved by modifying thecasting stacking pattern in the heat treatment baskets.

Figure 10. Photograph of an upper control arm casting in the machining fixture.

The close, open minded, co-operative workingrelationship between foundry and machiner personneland the proximity of the involved plants in Decatur, ILhave been major assets in implementing and launchingthe Ford Mustang ADI upper control arm program.

SUMMARY

This program has demonstrated that austempered ductileiron is a suitable contender for chassis applicationsrequiring lightweight, tough, high strength components.For the Ford Mustang upper control arm, it has provensuperior to a steel forging. With good process controlduring casting production and heat treatment and withsolid fixturing, ADI is readily machinable. It is clearly amaterial worthy of consideration for automotiveapplications

ACKNOWLEDGMENTS

We are grateful to all of the individuals at Ford, Benteler,Applied Process and Climate Control who supported usin this program. We would like to particularly recognizeCindy Jacob at Ford Motor Company for her numeroustechnical contributions to the program, and Monte Langeat Climate Control for his ongoing support, the time hespent with us in early phases of the development of thispaper, and for the insights he provided regarding themachining of ADI. Other individuals deserving special

recognition are Mitch Rackers for his handling of overallcomponent design and design analysis, and Steve Braunfor his development of the nondestructive test proceduresused to confirm successful austempering of eachcasting.

REFERENCES

1. SAE J434 Jun 86, 1998 SAE Handbook on CD-ROM,SAE International, Warrendale, PA.

2. ASTM A 536, 1999 Annual Book of ASTMStandards, Vol. 1.02, ASTM, West Conshohocken,PA, pp 310-314.

3. ASTM A897, 1999 Annual Book of ASTM Standards,Vol. 1.02, ASTM, West Conshohocken, PA, pp 570-575.

4. Natalie Neff, Mustang Cobra Gets Audited for 1999,WARD’S Auto World, February 1999, p 47

5. Norman Martin, Mustang Finally Gets An IRS, AI,March 1999, pp 63-64

6. Intermet internal studies.

CONTACT

Dr. Robert J Warrick received his degree in MetallurgicalEngineering from the University of Michigan in 1963. Hisindustrial career has been split between Ford MotorCompany (first 14 years) and Intermet. He is currentlyVice-President Materials R&D for Intermet Corporationand is located at the Intermet Product Design andTechnical Center, 939 Airport Road, Lynchburg VA24502. He can be reached by phone at (804) 237-8747or at rwarrick@notes.intermet.com by email.

AUSTEMPERING DUCTILE IRON

We want to present several of the ideas important to the successful austempering of ductile iron to help you decide if this is a market where you can participate. The ideas presented are generalizations - over simplifications - and we would welcome every opportunity to plan specific jobs with you.

APPLICATIONS: Austempered Ductile Iron (ADI) has been successfully applied to a variety of parts. Of great importance to the foundry industry is that much of it has replaced forgings. Specifics follow.

In-ground applications: Farm and construction equipment. Material handling equipment: Conveyor components, wheels, links. Gears: Automotive drive trains. Cam and Crank Shafts. Heavy duty suspension components: Spring hangers, equalizers.

ECONOMICS: ADI derives its value in any combination of these three areas: Superior performance in service. Casting cost advantage over other fabrication methods. Machining cost advantage of ductile iron over other materials.

TECHNOLOGY: ADI is produced in a range of strengths from 80,000 psi yield to 185,000 psi yield. As the strength increases through this range, the typical elongation drops from 13% to 2%. The following are important considerations:

The machinability of lower strength grades of ADI is reasonable -roughly comparable to steels of the same hardness. When hardness exceeds 321 BHN, parts are best ground. In most cases, it is best to machine parts prior to Austempering.

Dimensional stability is good since the ADI process is "Isothermal." Accordingly, cracking and warpage are rarely issues. Parts, however, do grow about 0.0008 in/in and very predictably, so machining dimensions are easily adjusted to compensate.

Metal chemistry needs to meet only a few criterion. Most important is for the chemistry to be consistent from lot to lot. Second in importance is manganese content which is best kept below 0.35%. Please note that this becomes increasingly important with increasing wall sections. Third is alloying and wall section considerations. To get a good ADI structure through walls greater that 0.75 inches, we usually need a little copper; or occasionally, nickel or molybdenum.

The Austenitizing time and temperature is determined by the metal chemistry.

The ADI strength is controlled by the Austemper time and temperature, and it is not significantly affected by metal chemistry. This means that we can obtain all grades of ADI with one chemistry of iron.

ASTM 897-90 FOR AUSTEMPERED DUCTILE IRON GRADE BHN

RANGE U.T.S. (PSI)

YIELD (PSI)

ELONG %

IMPACT FT/LBS

APPLICATIONS

1 269-321 125 80 10 75 Trunions, Hubs, Gears, Levers, Brackets

2 302-363 150 100 7 60 Chain links, Sprockets Crane Wheels

3 341-444 175 125 4 45 Wear plates, Mill liners Plow shoes, Axe heads, Mixer shoes

4 388-477 200 155 1 25 Rock crusher rolls, Chute liners, Rack and pinions, Guides

5 444-555 230 185 N/A N/A Break shoes and other high wear applications

Characteristics

Grades 1 through 3 have superior strength without the attending loss of ductility typical of ductile cast iron. Tensile and yield strengths are, in fact, comparable to low alloy steels.

Grades 4 and 5, with higher hardnesses, provide superior wear resistance and can provide cost effective substitutes for many carburized and quenched and tempered steel components.

It is possible, especially in the higher grades, to attain an ADI microstructure which contains some retained austenite. This phase will transform in service just at the working surface which further improves wear resistance.

The modulus of elasticity of ADI is 20% less than that of steel. This results in 40% faster dampening in ADI than in steel. When steel gears are replaced by ADI gears, noise reduction is typically five (5) decibels.

Ductile cast iron is about 10% lighter (less dense) than steel. Replacement of steel components with ADI will reduce weight.

Fatigue Performance Comparison and Life Predictions of Forged Steel and Ductile Cast Iron Crankshafts

Jonathan R. Williams and Ali Fatemi Graduate Research Assistant and Professor, Respectively

A Final Project Report Submitted to the

Forging Industry Educational Research Foundation (FIERF) and

American Iron and Steel Institute (AISI)

The University of Toledo

August 2007

iii

FORWARD

The overall objective of this study was to evaluate and compare the fatigue

performance of two competing manufacturing technologies for automotive crankshafts,

namely forged steel and ductile cast iron. In addition, weight and cost reduction

opportunities for optimization of the forged steel crankshaft were also investigated. The

detailed results are presented in two reports. This first report deals with the fatigue

performance and comparison of forged steel and ductile cast iron crankshafts. The

second report deals with analyses of weight and cost reduction for optimization of the

forged steel crankshaft.

iii

ABSTRACT

Fatigue Performance Comparison and Life Predictions of

Forged Steel and Ductile Cast Iron Crankshafts

The primary objective of this study was to evaluate and compare the fatigue

performance of forged steel and ductile cast iron crankshafts. Fatigue is the primary

cause of failure of crankshafts due to the cyclic loading and presence of stress

concentrations at the fillets. The crankshafts used in this study were from one-cylinder

engines typically used in lawnmowers. Recent publications relevant to this work are

presented. The experimental program included monotonic tensile tests, strain-controlled

fatigue tests, Charpy V-notch impact tests, as well as load-controlled component fatigue

tests on both crankshafts. Monotonic and cyclic properties of the two materials were

obtained and compared, which showed a higher tensile strength and better fatigue

performance for the forged steel compared to the ductile cast iron. The results from the

Charpy V-notch tests showed that the forged steel material has higher impact toughness

than the ductile cast iron material. The results of the component fatigue tests are

presented as S-N curves for the two crankshafts and show a superior fatigue performance

for the forged steel crankshafts. In addition to the experimental program, life predictions

were performed for the two crankshafts using the properties obtained from the strain-

controlled specimen tests. Results from FEA were used to determine the stress

concentrations in the crankshafts along with the stress distributions. S-N life predictions

were performed using the modified Goodman equation to account for the mean stress

iv

effects caused by the R = -0.2 loading. Strain-life predictions were also performed using

Neuber’s rule to determine the notch stresses and strains and the SWT parameter for

accounting for mean stress effects. Both the S-N and strain-life predictions provided

reasonable estimates to the fatigue lives of the crankshafts, although the S-N predictions

were in better agreement with the experimental data than the strain-life predictions.

v

ACKNOWLEDGEMENTS

Financial support for this research project was provided by the Forging Industry

Educational Research Foundation (FIERF) and the American Iron and Steel Institute

(AISI). We would like to thank Karen Lewis (Executive Director of FIERF), David

Anderson (Director of Bar and Rod Products at AISI), Michael Wicklund (President of

FIERF) for providing technical support and information, and George Mochnal from the

Forging Industry Association. In addition we would like to acknowledge Bill Heitmann

and Louie Laus of Arcelor Mittal Steel for there generous help and assistance with the

chemical analyses and microstructure imaging.

vi

TABLE OF CONTENTS FORWARD………………………………………………………………………………ii ABSTRACT……………………………………………………………………………...iii ACKNOWLEGEMENTS…………………………………………………………………v LIST OF TABLES…………………………………………………………….……...…..ix LIST OF FIGURES……………………………………………………………………....xi CHAPTER 1: INTRODUUCTION………………….……………………………………1 1.1 Background……………………………………………………………………………1

1.1.1 Crankshaft description………………………………………………….................1 1.1.2 Function of a crankshaft…………………………………………………………..4

1.1.3 Failure of a crankshaft…………………………………………………………….5

1.2 Literature Review……………………………………………………………………..6

1.2.1 Failure analysis……………………………………………………………………7

1.2.2 Testing and comparison of fatigue performance of crankshafts…………………..8

1.2.3 Crankshaft manufacturing………………………………………………………..11 1.3 Motivation and Objectives…………………………………………..……………….12 CHAPTER 2: SPECIMEN TESTING PROCEDURES AND RESULTS……………....26 2.1 Monotonic and Fatigue Tests and Results…………………………………………...26

2.1.1 Materials, specimen and test equipment…………………………………………26

2.1.2 Test procedures…………………………………………………………………..30

2.1.2.1 Monotonic tension tests………………………………………………….30

2.1.2.2 Constant amplitude fatigue tests…………………………………………31

2.1.3 Experimental results and comparisons…………………………………………...33

vii

2.1.3.1 Monotonic properties…………………………………………………….33 2.1.3.2 Cyclic deformation properties and behavior……………………………..35

2.1.3.3 Fatigue behavior and comparisons……………………………………….38

2.2 Charpy V-Notch Tests...……………………………………………………………...41

2.2.1 Specimen and test equipment…………………………………………………….41

2.2.2 Test procedure……………………………………………………………………43

2.2.3 Test results and comparisons…………………………………………………….44 CHAPTER 3: COMPONENT TESTING PROCEDURES AND RESULTS.…………..81 3.1 Test Apparatus and Procedures……………………………………………………....81

3.1.1 Loading conditions and test fixture………………………………………………81

3.1.2 Test procedures…………………………………………………………………...83 3.2 Failure Criterion……………………………………………………………………....84 3.3 Results and Comparisons ……………………………………………….……………88 CHAPTER 4: STRESS ANALYSIS AND FATIGUE LIFE PREDICTIONS………....107 4.1 Analytical Stress Calculations…………………………………………………….....107 4.2 Finite Element Modeling and Analysis...…………………..………………………...108

4.2.1 Critical locations……………………………………………………..…………..109

4.2.2 Comparison between FEA, analytical, and experimental results…………..……111

4.2.3 FEA Results used for life predictions………..…………………………...……...112 4.3 Stress-Life Approach and Life Predictions………………………………….…........114

4.3.1 Procedures and predictions.. …………………………………………………….114

4.3.2 Comparisons with experimental results..………………………………..…..…..119 4.4 Strain-Life Approach and Life Predictions.………………………………………....122

viii

4.4.1 Procedures and predictions...………………………………………………….…122

4.4.2 Comparisons with experimental results……...…………………………………125

4.5 Discussion of Life Prediction Results……………….………….…………………..126 CHAPTER 5: SUMMARY AND CONCLUSIONS…………………………………...142 REFERENCES…………………………………………………………………………146

ix

LIST OF TABLES

Table 1.1: Results from component fatigue tests on forged steel, ductile iron and

ADI crankshafts with various surface treatments from the study by Chatterley and Murrell [1998]. ....................................................................... 15

Table 1.2: Results from component fatigue tests on forged steel, ductile iron, and

microalloyed steel crankshafts from the study by Pichard et al. [1993]. ....... 15 Table 2.1: Chemical analysis of the forged steel and ductile cast iron as a percent

weight, remaining Fe [Heitmann, 2006]. ........................................................ 46 Table 2.2: Hardness values for (a) forged steel and (b) ductile cast iron

monotonic and fatigue specimens. .................................................................. 47 Table 2.3: Result summary of monotonic tensile tests. ................................................... 48 Table 2.4: Summary of monotonic and cyclic properties for the two materials.............. 49 Table 2.5: Summary of constant amplitude completely reversed fatigue test

results for forged steel..................................................................................... 50 Table 2.6: Summary of constant amplitude completely reversed fatigue test

results for ductile cast iron.............................................................................. 51 Table 2.7: Summary of results from Charpy impact tests for (a) forged steel L-T,

(b) forged steel T-L, and (c) cast iron. ............................................................ 52

Table 3.1: Test parameters and results for the forged steel and ductile cast iron crankshaft fatigue tests.................................................................................... 92

Table 4.1: Analytical nominal stress results at the critical location and

comparison with FEA results for the forged steel and cast iron crankshafts. ................................................................................................... 128

Table 4.2: Comparison between FEA, experimental, and analytical stress results

for the forged steel crankshaft in the as-tested condition at the locations shown in Figure 4.2. ...................................................................... 129

Table 4.3: FEA results for the test setup boundary conditions for the forged steel

crankshaft for the locations identified in Figure 4.2. .................................... 130

x

Table 4.4: FEA results for the test setup boundary conditions for the cast iron crankshaft for the locations identified in Figure 4.2. .................................... 130

Table 4.5: Life prediction results including the S-N and ε-N approaches for the

forged steel crankshaft. ................................................................................. 131 Table 4.6: Life prediction results including the S-N and ε-N approaches for the

ductile cast iron crankshaft. .......................................................................... 131 Table 4.7: Experimental data and life prediction results for the forged steel and

ductile cast iron crankshafts.......................................................................... 132

xi

LIST OF FIGURES

Figure 1.1: Crankshaft terminology [www.tpub.com]. ................................................... 16

Figure 1.2: The cycles of a four-stroke engine [en.wikipedia.org]. ................................ 16

Figure 1.3: Geometry of one cylinder diesel crankshaft used in the study by Bayrakçeken et al. [2006]. ............................................................................ 17

Figure 1.4: Fracture surfaces from failed one-cylinder diesel crankshafts from the study by Bayrakçeken et al. [2006]. ............................................................ 17

Figure 1.5: SEM photographs of failed crankshafts showing carbide inclusions indicated with arrows from the study by Bayrakçeken et al. [2006]. ........... 17

Figure 1.6: Failed crankshaft from a 6-cylinder diesel engine from the study by Asi [2006].................................................................................................... 18

Figure 1.7: Close up of crack in failed crankshaft from the study by Asi [2006]. .......... 18

Figure 1.8: Circumferential crack in failed crankshaft from the study by Asi [2006]. ........................................................................................................... 18

Figure 1.9: SEM photograph of crack initiation site in the fillet region from the study by Asi [2006]. ...................................................................................... 19

Figure 1.10: Test set-up to determine the modal response of specimens from the study by Damir et al. [2007]........................................................................ 19

Figure 1.11: Damping ratio versus life to failure for grey cast iron and ductile cast iron specimens from the study by Damir et al. [2007]................................ 20

Figure 1.12: Life to failure versus damping ratio for ADI specimens showing a quadratic correlation from the study by Damir et al. [2007]....................... 20

Figure 1.13: Test section for resonant bending test from the study by Spiteri et al. [2007]. ......................................................................................................... 21

Figure 1.14: Test apparatus for resonant bending fatigue test from the study by Spiteri et al [2007]....................................................................................... 21

Figure 1.15: Results from component tests on ductile cast iron crankshafts with various surface treatments from the study by Park et al. [2001]................. 22

xii

Figure 1.16: Electroslag casting (ESC) process shown schematically where A: transformer; B: Bottom mould; C1, C2, C3: mould; D: casting; E: molten metal pool; F: slag pool; G: electrode [Wang et al. 2007]. ............. 23

Figure 1.17: Forging sequence of the elementary cell for a precision forged crankshaft from the study by Behrens et al. [2005]. ................................... 23

Figure 1.18: Sequence for precision forging of a one-cylinder crankshaft from the study by Behrens et al. [2005]..................................................................... 24

Figure 1.19: Tool layout for the final forming stage of a one-cylinder crankshaft from the study by Behrens et al. [2005]. ..................................................... 24

Figure 1.20: Forging sequence for the precision forging of a three-cylinder crankshaft from the study by Behrens et al. [2005]. ................................... 25

Figure 2.1: Forged steel (a) and ductile cast iron (b) crankshafts used to obtain test specimens. ................................................................................................... 53

Figure 2.2: Photomicrographs of the ductile cast iron material at (a) 500X and (b) 1000X [Laus and Heitmann, 2007]. ............................................................ 54

Figure 2.3: Photomicrograph of the forged steel material at 500X................................. 55

Figure 2.4: Specimen geometry for monotonic tensile tests and constant amplitude fatigue tests.................................................................................................. 56

Figure 2.5: Locations where the monotonic and fatigue specimens were removed from for forged steel (a) and cast iron (b). .................................................. 57

Figure 2.6: True stress versus true plastic strain for (a) forged steel and (b) ductile cast iron. ...................................................................................................... 58

Figure 2.7: Monotonic engineering stress versus strain curves for (a) forged steel and (b) ductile cast iron. .............................................................................. 59

Figure 2.8: Superimposed monotonic engineering stress versus strain curves for forged steel and ductile cast iron................................................................. 60

Figure 2.9: True stress amplitude versus number of cycles for (a) forged steel and (b) ductile cast iron...................................................................................... 61

Figure 2.10: True stress amplitude versus normalized number of cycles for (a) forged steel and (b) ductile cast iron. .......................................................... 62

Figure 2.11: Plots of midlife hysteresis loops for (a) forged steel and (b) cast iron. ...... 63

Figure 2.12: True stress amplitude versus true plastic strain amplitude for (a) forged steel and (b) ductile cast iron. .......................................................... 64

xiii

Figure 2.13: True stress amplitude versus true strain amplitude for (a) forged steel and (b) ductile cast iron. .............................................................................. 65

Figure 2.14: Superimposed cyclic stress-strain curves for forged steel and ductile cast iron. ...................................................................................................... 66

Figure 2.15: Superimposed plots of monotonic and cyclic true stress versus true strain curves for (a) forged steel and (b) ductile cast iron........................... 67

Figure 2.16: Superimposed plots of monotonic and cyclic true stress versus true strain curves for forged steel and ductile cast iron...................................... 68

Figure 2.17: True stress amplitude versus reversals to failure for (a) forged steel and (b) ductile cast iron. .............................................................................. 69

Figure 2.18: Superimposed plots of true stress amplitude versus reversals to failure for forged steel and ductile cast iron. .......................................................... 70

Figure 2.19: True plastic strain amplitude versus reversals to failure for (a) forged steel and (b) ductile cast iron....................................................................... 71

Figure 2.20: Superimposed plots of true plastic strain versus reversals to failure for forged steel and ductile cast iron................................................................. 72

Figure 2.21: True strain amplitude versus reversals to failure for (a) forged steel and (b) ductile cast ...................................................................................... 73

Figure 2.22: True strain amplitude versus reversals to failure for forged steel and ductile cast iron. .......................................................................................... 74

Figure 2.23: Neuber stress range versus reversals to failure for (a) forged steel and (b) ductile cast iron...................................................................................... 75

Figure 2.24: Superimposed Neuber stress range versus reversals to failure for forged steel and ductile cast iron................................................................. 76

Figure 2.25: Charpy impact specimen geometry............................................................. 77

Figure 2.26: Locations on the crankshaft where Charpy v-notch specimens were machined from ...…………………………………………………………..77

Figure 2.27: Tinius Olsen impact test machine. .............................................................. 78

Figure 2.28: Average absorbed energy values at the different test temperatures for forged steel (L-T, T-L) and ductile cast iron............................................... 79

Figure 2.29: Absorbed energy versus test temperature for forged steel (L-T, T-L) and cast iron specimens............................................................................... 79

xiv

Figure 2.30: Fracture surfaces of (a) forged steel L-T, (b) forged steel T-L, and (c)

ductile cast iron specimens in order of increasing temperature from left to right………………………………………………………………...80

Figure 3.1: Forged steel crankshaft in its final machined condition. .............................. 93

Figure 3.2: Ductile cast iron crankshaft in its final machined condition. ....................... 93

Figure 3.3: Schematic of test set-up. ............................................................................... 94

Figure 3.4: Test set-up for the forged steel crankshaft.................................................... 94

Figure 3.5: Test set-up for the ductile cast iron crankshaft. ............................................ 95

Figure 3.6: Close up of load application area of moment arm showing rod end bearing and roller bearings. ......................................................................... 95

Figure 3.7: Critical fillet area of crankshaft painted to better observe crack. ................. 96

Figure 3.8: Imprint of crack with putty. .......................................................................... 96

Figure 3.9: Displacement amplitude versus number of cycles for the (a) forged steel crankshafts and (b) ductile cast iron crankshafts. ............................... 97

Figure 3.10: Change in displacement amplitude versus crack length for the forged steel crankshafts. ......................................................................................... 98

Figure 3.11: Change in displacement amplitude versus crack length for the cast iron crankshafts. .......................................................................................... 98

Figure 3.12: Superimposed plot of change in displacement amplitude versus crack length for the forged steel and cast iron crankshafts. ................................. 99

Figure 3.13: Measured crack length versus cycles for the forged steel crankshafts. .... 100

Figure 3.14: Measured crack length versus cycles for the ductile cast iron crankshafts................................................................................................. 100

Figure 3.15: Displacement amplitude versus cycles for a forged steel crankshaft with the crack initiation point highlighted. ............................................... 101

Figure 3.16: Displacement amplitude versus cycles for a ductile cast iron crankshaft with the crack initiation point highlighted............................... 101

xv

Figure 3.17: Predicted crack length versus measured crack length for the forged steel crankshafts. The same symbols correspond to crack lengths of the same crankshaft. .................................................................................. 102

Figure 3.18: Predicted crack length versus measured crack length for the ductile cast iron crankshafts. The same symbols correspond to crack lengths of the same crankshaft. ............................................................................. 102

Figure 3.19: Moment amplitude versus cycles to failure using the crack initiation failure criterion. ......................................................................................... 103

Figure 3.20: Moment amplitude versus cycles to failure using the 5% change in displacement amplitude failure criterion................................................... 103

Figure 3.21: Cast Iron displacement amplitude versus cycles plot showing hardening behavior. ................................................................................... 104

Figure 3.22: Expanded view of the displacement amplitude versus cycles plot for a cast iron crankshaft tested at 431 N-m. ..................................................... 104

Figure 3.23: Example of a typical fatigue fracture surface for the forged steel crankshaft. ................................................................................................. 105

Figure 3.24: Side view of typical fatigue fractured forged steel crankshaft. ................ 105

Figure 3.25: Example of a typical fatigue fracture surface for the cast iron crankshaft. ................................................................................................. 106

Figure 3.26: Side view of typical fatigue fractured cast iron crankshaft....................... 106

Figure 4.1: Forged steel crankshaft showing FEA stress contour with the crankpin fillet magnified [Montazersadgh, 2007]...………………………………..133

Figure 4.2: Forged steel crankshaft showing the analyzed locations for the

dynamic load analysis and dynamic based FEA. ...................................... 133

Figure 4.3: Stress magnitude versus crankshaft angle for the locations shown in Figure 4.2 [Montazersadgh and Fatemi, 2007]. ........................................ 134

Figure 4.4: Maximum stress, minimum stress, stress range, and mean stress results from FEA for the locations shown in Figure 4.2 [Montazersadgh and Fatemi, 2007]. ........................................................................................... 134

Figure 4.5: Forged steel crankshaft S-N lines for the unnotched, notched, and notched ...................................................................................................... 135

Figure 4.6: Ductile cast iron crankshaft S-N lines for the unnotched, notched, and notched R = -0.2 condition. ....................................................................... 135

xvi

Figure 4.7: Forged steel crankshaft S-N line for the notched R = -0.2 condition superimposed with the crack initiation experimental data. ....................... 136

Figure 4.8: Forged steel crankshaft S-N line for the notched R = -0.2 condition superimposed with the 5% change in displacement amplitude experimental data. ..................................................................................... 136

Figure 4.9: Ductile cast iron crankshaft S-N lines for the notched R = -0.2 condition superimposed with the crack initiation experimental data. ....... 137

Figure 4.10: Ductile cast iron crankshaft S-N lines for the notched R = -0.2 condition superimposed with the 5% change in displacement amplitude experimental data. .................................................................... 137

Figure 4.11: Predicted versus experimental cycles to failure using the S-N approach for the forged steel and ductile cast iron crankshafts using the crack initiation failure criterion. .......................................................... 138

Figure 4.12: Predicted versus experimental cycles to failure using the S-N approach for the forged steel and ductile cast iron crankshafts using the 5% change in displacement amplitude failure criterion. ..................... 138

Figure 4.13: SWT parameter versus reversals to failure based on crack initiation with strain-life prediction data superimposed for the forged steel crankshafts................................................................................................. 139

Figure 4.14: SWT parameter versus reversals to failure based on 5% change in displacement amplitude with strain-life prediction data superimposed for the forged steel crankshafts. ................................................................ 139

Figure 4.15: SWT parameter versus reversals to failure based on crack initiation with strain-life prediction data superimposed for the ductile cast iron crankshafts................................................................................................. 140

Figure 4.16: SWT parameter versus reversals to failure based on 5% change in displacement amplitude with strain-life prediction data superimposed for the ductile cast iron crankshafts........................................................... 140

Figure 4.17: Predicted versus experimental cycles to failure using the strain-life approach for the forged steel and ductile cast iron crankshafts based on the crack initiation failure criterion. ..................................................... 141

Figure 4.18: Predicted versus experimental cycles to failure using the strain-life approach for the forged steel and ductile cast iron crankshafts based on the 5% change in displacement amplitude failure criterion. ................ 141

1

CHAPTER 1

INTRODUCTION

1.1 Background

A crankshaft, in general, converts linear motion into rotary motion. In an internal

combustion engine, the reciprocating motion of the piston is linear and is converted into

rotary motion through the crankshaft. The most common application of a crankshaft is in

an automobile engine. However, there are many other applications of a crankshaft which

range from small one cylinder lawnmower engines to very large multi cylinder marine

crankshafts and everything in between.

1.1.1 Crankshaft description

A crankshaft consists of main journals, webs, and connecting rod journals,

commonly known as “crank-pins”. The main components of a crankshaft are shown in

Figure 1.1. The crankshaft rotates on bearings inside the engine. The bearings

supporting the crankshaft are the main bearings of an engine and the part of the

crankshaft that rides on the bearings are called the main bearing journals. The number of

main bearings and main journals in an engine depend on its size. Small one cylinder

engines have only two main bearings, one at each end of the crankshaft. Larger multi-

cylinder engines usually have more than two main bearings at the ends and include some

in the center part of the crankshaft for more support. The piston connects to the

2

crankshaft on a bearing journal, referred to as a crank-pin. The crank-pins are offset from

the central rotating axis of the crankshaft causing the pistons to move when the

crankshaft rotates. The webs create the offset between the central axis and the crank-

pins. The number of crank-pins depends on the type of engine and number of cylinders.

A single –cylinder engine will have only one crank-pin and two webs. Multi-cylinder

engines will have one crank-pin per piston if the engine is a straight engine, meaning that

all cylinders are in a line. If the engine is a V-engine, one bank of cylinders on each side

of the crankshaft, two pistons will attach to the same crank-pin. Commonly a crankshaft

will be classified by the number of “crank throws” or simply “throws”, which simply

refers to the combination of the two webs and crank-pin. Therefore, a straight four

cylinder engine will have four crank-pins as will a V-8 engine and both will be classified

as four throw crankshafts.

The high speed rotation of a crankshaft requires a very balanced component. If

the crankshaft is not balanced damage to the engine can result or at the very least there

will be a heavy vibration. Balancing of a crankshaft is partly achieved by using counter

balance weights on the crankshaft. The webs usually extend past the central axis of the

crankshaft to form the counterweights. Fine balancing is usually done by drilling holes in

the underside of the counterweights to remove material. The locations of the holes are

such that when the material is removed the crankshaft will be in balance.

There are several different material options available for manufacturing

crankshafts, with the two most popular being steel and iron. Crankshafts can be

machined from a billet, forged, or cast. Machining a crankshaft from a billet is not

typically done due to the prohibitively long machining times, however for low production

3

custom pieces it is still done. The steel crankshaft is usually forged to near net shape and

then finished by machining processes. The iron crankshafts are typically cast and then

machined. Generally a crankshaft can be classified as forged steel or cast iron, however,

within these two categories there are many options. A forged steel crankshaft, for

example, may be manufactured from microalloyed steel which can eliminate the need for

heat treatment. A cast iron crankshaft, which is typically ductile cast iron, has more

ductility and therefore higher fatigue resistance than ordinary gray iron. The cast iron

crankshaft could also be made from austempered ductile iron or ADI, which is a higher

strength iron and has been shown to have longer fatigue life than ordinary ductile cast

iron [Chatterley and Murrell, 1998].

Crankshaft design is not limited to selecting a material, such as steel or iron, a

process, such as forging or casting, and geometry. Surface treatments also play a major

role in the performance of the crankshaft. The fillets in a crankshaft are often rolled in

order to induce compressive residual stresses, thus increasing the fatigue performance of

the crankshaft. Case hardening, or hardening on the surface of the material, is often done

to increase the hardness in the main journals and crank-pins of the crankshaft, resulting in

better wear. Not only does the surface hardening improve wear resistance, it also can

induce compressive residual stresses, which results in increased fatigue performance of

the crankshaft [Grum, 2003]. Ion nitriding is also used and has been shown to increase

the fatigue strength of crankshafts [Park et al., 2001; Pichard et al, 1993].

4

1.1.2 Function of a crankshaft

As mentioned previously, the function of a crankshaft in an internal combustion

engine is to translate the linear motion of the pistons into a rotary motion. The rotary

motion can then be used to power the device, such as propel an automobile or turn the

blade of a lawnmower engine. The most common application is the automobile engine.

The function of the crankshaft in an automobile engine can easily be extended to a

crankshaft in another type of engine since their functions are the same. Since the

crankshafts used in this study are from a four-cycle engine, the four-cycle engine process

is discussed. Crankshafts, however, perform similar functions in the two-cycle engine

which is common in small garden equipment.

The most common type of engine is the four-cycle (or four-stroke) engine which

uses the Otto cycle (if gasoline powered) or Diesel cycle (if diesel powered). The four

cycles are the intake, compression, power, and exhaust, which are shown in Figure 1.2.

The cycle starts at top dead center (TDC) where the piston is furthest away from the

crankshaft. In the first cycle, the intake cycle, the piston moves down and an air-fuel

mixture is drawn into the cylinder through the intake valves. Next the valves close and

the piston moves up, compressing the air-fuel mixture in the compression cycle. The

compressed air-fuel mixture is ignited (by a sparkplug in a gasoline engine) at the top of

the compression stroke. The power cycle occurs when the gases in the combustion

chamber ignite, resulting in expansion and a large force on the piston. The force pushes

the piston down resulting in a rotation of the crankshaft. Finally, in the exhaust cycle, the

exhaust valves open and the gases in the cylinder are forced out during the upward

5

motion of the piston as the crankshaft rotates. The entire process results in a 720 degree

rotation of the crankshaft, as each cycle takes approximately 180 degrees to complete.

1.1.3 Failure of a crankshaft

The crankshaft is the central part of the engine and its failure would render the

engine useless until costly repairs could be made or a replacement engine could be

installed. The failure of a crankshaft can damage other engine components including the

connecting rods or even the engine block itself. Therefore, when the failure of a

crankshaft does occur it often results in replacing the engine or even scrapping the

equipment the engine was used in. Considering the ramifications of a crankshaft failure,

a crankshaft must be designed to last the lifetime of an engine.

The engine of a typical gasoline powered automobile has an engine speed that

varies from 500 to 6,500 rpm and while traveling at highway speeds may be 2,500 rpm.

It can easily be shown that a crankshaft has a desired life of many millions or even

billions of cycles. For example if the life of an automobile is 120,000 miles and has an

average speed of 50 mph and engine speed of 2,500 rpm, the engine, and crankshaft,

would need to have a life of at 360 million cycles. Crankshafts used in lawnmower

engines, such as those in this study, would not see as many cycles as the automobile

engine with considerably higher usage, but a long life situation still exists. With such a

long life situation, a design for infinite life is necessary.

The gas and inertial loads in an engine create a multiaxial loading condition on a

crankshaft as was shown by Jensen [1970]. In the study strain gages were mounted to a

crankshaft from a V-8 engine and installed back in the engine. By running the engine

6

and acquiring data, he was able to show that there was bending and torsion on the

crankshaft. The study by Jensen and subsequent studies show that the torsion is small

compared to the bending stress, therefore, the torsion is often neglected.

The fillets in crankshafts have been identified as the highest stressed, or critical,

location of a crankshaft and are often the sight of fatigue crack initiation as was shown in

the previously mentioned study by Jensen and other studies, including this one. The

presence of a fillet or notch in a crankshaft is virtually unavoidable. Any change of

diameter results in a stress concentration. While sharp corners can be avoided with the

use of fillets, other measures are often necessary in order to increase the fatigue

performance of crankshafts. Compressive residual stresses have been shown to increase

the fatigue performance of components, not just limited to crankshafts. Often in an

attempt to induce compressive residual stresses at notches, the fillets are rolled. This

compressive residual stress increases the fatigue strength at long life.

Silva [2003] classified the failure of crankshafts into three categories: operating

sources, mechanical sources, and repairing sources. Operating sources include things

such as misuse of an engine and a lack of lubrication. Mechanical sources of failure can

include misalignment or vibration of the crankshaft due to balance issues. Repairing

sources are those that are caused by repair to an engine or finishing of a crankshaft, such

as improper grinding, incorrect bearings, or misalignment.

1.2 Literature Review

The literature review for this project by Zoroufi and Fatemi [2005] was completed

previously. Therefore the literature review included in this chapter only contains

7

additional information that was published after the previous literature review was

completed and information that is mentioned again due to specific application to this

study.

1.2.1 Failure analysis

The analysis of failed in service crankshafts is vital to laboratory crankshaft

studies, as it allows the researcher to better adapt experiments to real life situations as

well as validates results. Crankshaft studies, including this one, suggest that crankshaft

failures often occur in the crank-pin fillet areas, which is also supported by the analysis of

failed in service crankshafts.

Bayrakçeken et al. [2006] investigated the failure of a small one-cylinder diesel

engine used in agricultural applications. The analysis was performed on two crankshafts

made of AISI 4140 steel, one of which was case hardened. The basic crankshaft

geometry used in the study is shown in Figure 1.3. Both failures were attributed to

fatigue crack propagation in the fillet of the crankshafts. Fracture surfaces which show

typical fatigue fractures, as indicated by beach marks, from both crankshafts are shown in

Figure 1.4. The premature failure was suspected to be caused by the larger than normal

carbide inclusions present in the material as shown by the scanning electron microscope

(SEM) images in Figure 1.5.

Asi [2006] investigated the failure of a diesel crankshaft made of ductile cast iron.

The crankshaft was taken from a 6 cylinder 115 HP engine which is shown in Figures 1.6

and 1.7. The failure of the crankshaft resulted in “catastrophic failure of the engine” after

only 400 hours of service. Circumferential cracks were found by visual inspection in the

8

crankpin fillet region as shown in Figure 1.8. From high magnification analysis of the

fracture surfaces using SEM, the initiation of the cracks was at the surface in the region

of the fillet as shown in Figure 1.9. The free graphite and nonspheroidal graphite in

ductile cast iron act as notches, and therefore stress concentrations, and are often the

source of fatigue crack initiation. The initiation site of the cracks in the analyzed

crankshaft is in the periphery of graphite. The analysis showed that the failure occurred

due to rotating bending fatigue. The initiation of cracks which ultimately led to the

fracture of the crankshaft occurred in the crankpin fillet region.

1.2.2 Testing and comparison of fatigue performance of crankshafts

Fatigue testing typically requires destructive testing of both specimens and

components in order to characterize the fatigue performance of a material or compare two

materials. Damir et al. [2007], however, describe a process for nondestructive

comparisons of fatigue behavior using modal analysis. Dynamic impact testing on

simple, cylindrical geometry, was performed along with rotating bending fatigue tests on

the same material using standard specimen geometry. An impact hammer equipped with

a force transducer was used to excite the specimen, while an accelerometer was used to

measure the response. The test set-up used is shown in Figure 1.10. Ductile cast iron

and gray cast iron were used in the study. The damping ratio and fatigue life were

affected by the microstructure of the material. Within a family of materials, a trend was

observed between the damping ratio and fatigue life as shown in Figure 1.11. A higher

damping ratio indicated a higher fatigue life. Ductile (nodular) cast iron specimens

having a higher damping ratio also had a higher fatigue life than the gray cast iron

9

specimens. There was no trend observed between natural frequency or magnitude of the

frequency response function (FRF) and fatigue life. Using austempered ductile iron

(ADI) a quadratic relationship between damping ratio was developed as shown in Figure

1.12. The equation of the curve shown in Figure 1.12 can be used to predict the fatigue

life based on the damping ratio for ADI at a stress of 500 MPa.

Spiteri et al. [2007] experimentally and analytically investigated the fatigue

performance of a ductile cast iron crankshaft subjected to bending loads. The objective

of the study was to compare different failure criteria. Tests were performed on a sample

cut from the crankshafts that consisted of two main bearings, one crankpin, and two

webs, as shown in Figure 1.13. Resonant bending fatigue tests were performed on the

test samples such that the crankpin fillet area was the highest stressed location. The test

fixture and setup is shown in Figure 1.14 and was the same fixture used by Chien et al.

[2005]. The data for the surface crack initiation failure criterion was taken from the study

by Chien et al. and the fatigue limit using this criterion was found to be 414 N-m.

Testing was done to compare the resonant frequency drop with the size of the crack.

Using a frequency drop failure criterion, the fatigue limit was found to be 642 N-m.

Therefore, the surface crack failure criterion was lower than the fracture criterion.

Papers by Park et al. [2001], Chatterley and Murrell [1998], and Pichard et al.

[1993] are included in the previous literature review [Zoroufi and Fatemi, 2005],

however, due to their direct relevance to this study they are mentioned here again.

In a study by Park et al. [2001] the effect of surface modifications was studied on

microalloyed CrMo crankshafts. The effect of fillet rolling using different forces as well

as nitriding was investigated. The results from component fatigue tests on the materials

10

and treatments are shown in Figure 1.15. It can be seen that many of the tests were

conducted to yield lives between 105 and 106 cycles, which compare to the types of tests

conducted in this study. The results show that a higher fillet rolling force induces higher

compressive residual stresses and in turn, produces better fatigue strength. However,

forces too high can be detrimental to the fatigue performance. Fillet rolling and nitrided

samples both produced approximately a factor of 1.8 increase in fatigue limit when

compared to untreated samples.

Chatterley and Murrell [1998] compared the fatigue performance of several

materials for use in a four-cylinder turbo charged diesel engine. The materials tested in

the study were nitrided 1% CrMo, fillet rolled ductile iron (Su = 700MPa), and fillet

rolled austempered ductile iron (ADI). Constant amplitude bending fatigue tests were

conducted on the crankshafts to 107 cycles or failure, whichever occurred first. The

results from the study are summarized in Table 1.1. The results indicate that nitrided

forged steel had a higher fatigue strength (107 cycles) than ductile iron or ADI regardless

of their surface treatments. The results also showed that ADI with higher rolling forces

had higher fatigue strength than the rolled ductile iron.

The study by Pichard et al. [1993] explored the possibility of replacing forged

steel or cast iron with a mircoalloyed steel in order to eliminate the need for additional

heat treatments. Tests on ductile cast iron, 1042 steel, 35MV7 microalloyed steel, and

32CDV13 high alloyed steel crankshafts were conducted. The results from the

component tests are shown in Table 1.2. The results showed that the control cooled

microalloyed steel had a higher fatigue strength than the 1042 steel and the ductile iron

for short nitriding treatments. The quenched and tempered 1042 steel did show

11

significantly higher fatigue strength than the ductile iron with the same surface treatment.

The microalloyed 35MV7 nitrided for 4 hours had only 10% lower fatigue strength than

the high alloyed 32CDV13 steel nitrided for 7 hours. This slightly lower fatigue strength

was combined with a significant cost savings, making the microalloyed steel a significant

contender.

1.2.3 Crankshaft manufacturing

As mentioned previously, there are several options available for manufacturing

crankshafts, most commonly casting and forging. While casting and forging are

generally used for high volume, ordinary sized crankshafts, alternative processes have

been investigated for low volume very large scale crankshafts. Wang et al. [2007]

discuss the fabrication of a large scale locomotive crankshaft using the electro-slag

casting (ESC) process which is shown schematically in Figure 1.16. Each manufacturing

process has its positive and negative attributes. These processes compete against one

another for strength, cost, efficiency, and production time motivations.

In order to decrease the amount of machining time, a precision forging technique

is used to produce forgings that are near net shape. Precision forging is a flashless

forging operation and has been used to produce small pieces such as connecting rods or

hand tools. Behrens et al. [2005] discuss the process of extending the precision forging

technique to larger, more complex shapes such as the crankshaft. Precision forging is a

hot forging process that uses closed dies. The process consists of an upper and lower die

and one or more punches. In order to develop a process for the complex crankshaft

shape, a series of steps were used. First the elementary cell, consisting of one main

12

bearing, one web, and one crankpin, was developed and verified which is shown in

Figure 1.17. After successful results with the elementary cell, the process was extended

to the one-cylinder crankshaft, consisting of one crankpin, two webs, and two half main

bearings (one on each side). The forging sequence for the one-cylinder crankshaft is

shown in Figure 1.18 with the final stage tool setup shown in Figure 1.19. Finally the

procedure for a three-cylinder crankshaft was developed. The forging sequence for the

three-cylinder crankshaft is shown in Figure 1.20. In the first process the main bearings

and the crankpins are formed. In the second process the webs are compressed and the

crankpins are translated to their eccentric position. The final step involves a tool with

punches integrated into the top and bottom dies. The dies in the final step represent the

shape of the main bearings and the crankpins that were formed in previous steps. The

punches in the final step form the geometry of the web sections.

1.3 Motivation and Objectives

In many industries, especially the automotive industry, there is a constant demand

for components that have less mass, are stronger, and cost less to produce. The

automotive industry, in particular, often seeks to improve gas mileage by using lighter

components, including optimized geometry and materials, all while reducing the cost of

manufacturing. Because of this, there is a constant debate over which material and

manufacturing process can be the most cost effective and lightest weight without

sacrificing performance.

The objective of this study was to assess and compare the fatigue performance of

forged steel and ductile cast iron crankshafts from a one-cylinder engine typical to that

13

used in a riding lawnmower. The forged steel crankshaft was designed to be used in a

460cc engine which produces approximately 9.3 kW. The ductile cast iron crankshaft

was from a similar engine size and type. The masses of both crankshafts were similar

with the forged steel at 3.9 kg and the ductile cast iron at 3.7 kg. Fatigue and monotonic

tests were conducted on standard specimens machined from the forged steel and ductile

cast iron crankshafts to compare the two materials. Component tests on both crankshafts

were conducted to obtain the fatigue properties and compare the two crankshafts. Finite

element analysis was used to determine the critical location of the crankshafts and to

determine the stress concentration factors. Life predictions were performed using both

the S-N approach and the strain-life approach, results of which were compared with the

component test data.

Dynamic load and stress analysis on the forged steel and ductile cast iron

crankshafts used in this study, as well as optimization of the forged steel crankshaft was

performed in another study, details of which can be found in Montazersadgh [2007].

Other details of the crankshaft study presented in this work can also be found in Zoroufi

and Fatemi [2005], Williams and Fatemi [2007], Montazersadgh and Fatemi, [2007], and

Williams et al. [2007].

The crankshafts used, being from a one-cylinder engine, were single throw

crankshaft consisting of two web sections and a one crankpin. Typically in automotive

crankshaft analysis a single throw is analyzed regardless of the size of the crankshaft.

Literature shows that for automotive crankshaft component fatigue tests, the crankshafts

are often sectioned so that a single throw can be tested [Chien et al., 2005; Spiteri et al.,

2007]. Therefore, the analyzed section in automotive crankshafts closely resembles the

14

analyzed section in this study, allowing the procedures and information to be easily

applied to automotive applications. Also, the failure location of the crankshafts used in

this study was in the crank-pin fillet, which agrees with the typical failure location for an

automotive crankshaft [Jensen, 1970].

Chapter 2 provides detailed description of the test procedures, results, and

comparisons of the specimen monotonic and fatigue tests as well as Charpy V-notch

impact tests. The Charpy V-notch tests were conducted due to the possibility that the

lawnmower contacts a hard object causing the engine to stop suddenly, thus resulting in

an impact loading condition in the engine. Chapter 3 describes the component fatigue

test procedures, results, and comparisons. The results of the stress analysis and FEA

performed are then discussed in Chapter 4. Chapter 4 also describes the life predictions

used and compares the results of the component tests with the life predictions. Finally,

Chapter 5 summarizes the conclusions made from the study.

15

Table 1.1: Results from component fatigue tests on forged steel, ductile iron and ADI

crankshafts with various surface treatments from the study by Chatterley and Murrell [1998].

Table 1.2: Results from component fatigue tests on forged steel, ductile iron, and microalloyed steel crankshafts from the study by Pichard et al. [1993].

16

Figure 1.1: Crankshaft terminology [www.tpub.com].

Figure 1.2: The cycles of a four-stroke engine [en.wikipedia.org].

17

Figure 1.3: Geometry of one cylinder diesel crankshaft used in the study by Bayrakçeken et al. [2006].

Figure 1.4: Fracture surfaces from failed one-cylinder diesel crankshafts from the study by Bayrakçeken et al. [2006].

Figure 1.5: SEM photographs of failed crankshafts showing carbide inclusions indicated with arrows from the study by Bayrakçeken et al. [2006].

18

Figure 1.6: Failed crankshaft from a 6-cylinder diesel engine from the study by Asi [2006].

Figure 1.7: Close up of crack in failed crankshaft from the study by Asi [2006].

Figure 1.8: Circumferential crack in failed crankshaft from the study by Asi [2006].

19

Figure 1.9: SEM photograph of crack initiation site in the fillet region from the study by Asi [2006].

Figure 1.10: Test set-up to determine the modal response of specimens from the study by Damir et al. [2007].

20

Figure 1.11: Damping ratio versus life to failure for grey cast iron and ductile cast iron specimens from the study by Damir et al. [2007].

Figure 1.12: Life to failure versus damping ratio for ADI specimens showing a quadratic correlation from the study by Damir et al. [2007].

21

Figure 1.13: Test section for resonant bending test from the study by Spiteri et al. [2007].

Figure 1.14: Test apparatus for resonant bending fatigue test from the study by Spiteri et al [2007].

22

Figure 1.15: Results from component tests on ductile cast iron crankshafts with various surface treatments from the study by Park et al. [2001].

23

Figure 1.16: Electroslag casting (ESC) process shown schematically where A: transformer; B: Bottom mould; C1, C2, C3: mould; D: casting; E: molten metal pool; F: slag pool; G: electrode [Wang et al. 2007].

Figure 1.17: Forging sequence of the elementary cell for a precision forged crankshaft

from the study by Behrens et al. [2005].

24

Figure 1.18: Sequence for precision forging of a one-cylinder crankshaft from the study

by Behrens et al. [2005].

Figure 1.19: Tool layout for the final forming stage of a one-cylinder crankshaft from the

study by Behrens et al. [2005].

25

Figure 1.20: Forging sequence for the precision forging of a three-cylinder crankshaft

from the study by Behrens et al. [2005].

26

CHAPTER 2

SPECIMEN TESTING PROCEDURES AND RESULTS

2.1 Monotonic and Fatigue Tests and Results 2.1.1 Materials, specimen, and test equipment

Ductile cast iron crankshafts and 1045 forged steel crankshafts were used in this

study. Both crankshafts were designed to be used in a one-cylinder small engine typical

to those found in riding lawn mowers. The crankshafts used for obtaining test specimens

were in the as cast and as forged condition when received and had not yet been through

the final machining process. The ductile iron and forged steel crankshafts in their

unmachined state are shown in Figure 2.1.

While limited information was available to identify the exact materials used in the

crankshaft, chemical analysis confirmed that the forged steel crankshaft was AISI 1045

steel. The chemical analysis along with microstructure analysis also confirmed that the

cast iron crankshaft was in fact a ductile cast iron. The chemical composition was

obtained from small sections removed from the as cast and as forged crankshafts. The

results from the chemical analyses of samples taken from the forged steel and ductile cast

iron are given in Table 2.1.

27

The microstructure of the ductile cast iron material consisted of spheriodal

graphite particles surrounded by patches of ferrite in a pearlite matrix. The optical

photomicrograph at 500X is shown for the ductile cast iron in Figure 2.2a. The scanning

electron microscope (SEM) photomicrograph at 1000X for the ductile cast iron is shown

in Figure 2.2b. The microstructure of the forged steel material was ferrite-pearlite. An

optical photomicrograph at 500X is shown in Figure 2.3.

Round specimens having the dimensions shown in Figure 2.4 were machined

from the two materials. Two specimens were machined from each crankshaft. The

longitudinal axis of the specimens coincided with the longitudinal axis of the crankshafts.

The locations where the specimens were taken from are shown in Figure 2.5. The

specimen geometry was slightly modified from the ASTM Standard E606 [2004]. The

standard specifies uniform, or hourglass specimens, while the specimen geometry used in

this study has a large secondary radius in the test section. The length of the grip section

was also shortened such that the specimens could be taken from the crankshafts, which

had limited usable length. Machining was performed by The University of Toledo

Mechanical, Industrial, and Manufacturing Engineering Machine Shop. The specimens

were rough cut to an approximate length from the indicated locations and then turned

down on a lathe to the desired diameter of the grip section. They were then cut to length

and center drilled. The final dimensions were machined using a CNC lathe.

The gauge section of each specimen was polished to remove all machining marks

by fixing one end of the specimen in the lathe and running the machine with a speed of

720 rpm. The polishing was done in five steps starting with the coarsest and ending with

the finest. Four of the steps were accomplished with small strips of sandpaper with grits

28

of 400, 600, 800, and 1000. The paper was kept wet by dipping it in water repeatedly

during the polishing process. The 400 grit paper was used until all machining marks

were removed. The subsequent steps were performed to remove the marks left by the

previous polishing process. The final of the five steps was a high speed rotating die

grinder with a buffing wheel attached. A polishing compound was applied to the wheel

and then the wheel was placed on the rotating specimen such that the direction of the

wheel was along the longitudinal axis of the specimen so that any marks from the final

step would be in the direction of the applied load during testing, minimizing their effect.

The specimens were carefully examined prior to testing to ensure that all marks were

removed from the test section.

Testing was performed on an Instron 8801 closed loop servo-hydraulic axial load

frame in combination with a Fast Track 8800 digital servo-controller. The load frame

was fitted with a 50 kN capacity load cell. The calibration of the system was verified

prior to testing. Collet type hydraulic grips were used in the test program. To insure that

the grips would maintain proper gripping of the specimen that had a grip section shorter

than the standard type specimen, spacer blocks were machined to fit in the collets of the

grip at each end of the specimen.

For the tests conducted in strain control an Epsilon Extensometer Model 3442 was

used to control the total strain. The extensometer conformed to ASTM Standard E83

[2004]. A verification procedure was performed on the extensometer to ensure proper

calibration. A displacement apparatus with a micrometer head (smallest increment of

measure 0.0001 inch) was used for the verification. The gauge length of the extensometer

was 6 mm (0.02362 in) and had a range of -6% to 10%. Each specimen was coated with

29

M-coat D at the locations of the extensometer edges prior to each test. This coating was

applied in order to prevent the knife edges of the extensometer from causing damage to

the specimen in the form of a stress concentration that could lead to fatigue failure at that

point. Prior to testing the extensometer was installed at the center of the specimens gauge

section and special care was taken to ensure that the extensometer was oriented parallel to

the load direction. Prior to performing a test the extensometer was allowed to reach a

state of stability by allowing it to be attached to the specimen for approximately one hour

(or until the changes in extensometer readout had stabilized). The tests were not

conducted until the reading of the extensometer was stable.

The environment in which the tests were carried out was carefully monitored and

maintained in order to minimize the effects on the extensometer and load cell due to

temperature. The ambient temperature was monitored during testing. The relative

humidity was also monitored using a hydrometer.

Due to the fact that any misalignment in the load train can result in inaccurate

tests caused by bending in the specimen, particular care was used in achieving proper

alignment. A precision round bar was fitted with eight strain gauges and placed in the

grips. The fine alignment adjustments were made with the Instron alignment fixture.

This fixture allows alignment adjustment that can result from tilt and/or eccentricity

between the central axis of the load train. Satisfactory alignment was achieved when the

bending was less than 5% of the axial load throughout the entire loading range used for

the testing. This is in accordance with the ASTM Standard E606 [2004] which prohibits

bending strains greater than 5% of the minimum axial strain range used for any fatigue

test.

30

2.1.2 Test procedures

2.1.2.1 Monotonic tension tests

A monotonic tension test was performed on each of the two materials. The test

was conducted in accordance with ASTM Standard E8 [2004]. Prior to testing the

location of the extensometer edges was marked on the specimen and the diameter of the

gauge section was measured using a 10X magnification optical comparator. The

software recorded the stress and strain data during the test in order to generate a stress-

strain curve. The extensometer has a maximum strain of 10% and the forged steel test

would have exceeded this value. The forged steel test was stopped prior to reaching 10%

strain and continued in displacement control until fracture. The cast iron specimen did

not approach the maximum range of the extensometer.

A strain rate of 0.0025 mm/mm/min was used from 0% to 0.5% strain. This

region represents the elastic and initial yield portion of the curve. This strain rate was

chosen because it is 75% of the maximum allowable strain rate specified by ASTM

Standard E8 [2004] for the initial yield region. After the specimen yielded, from 0.5% to

10% strain, the strain rate was increased by a factor of three to 0.0075 mm/mm/min. For

the forged steel test, once it reached near 10% strain (maximum permissible due to the

limitations of the extensometer) the test was switched to displacement control and a rate

of 0.152 mm/min was used. The rate of displacement was chosen to approximate the rate

of strain that the material experienced when the extensometer was attached.

Following the conclusion of the tests, the specimens were reassembled to their

prior to tested state. The final gauge length of the specimen was measured using a digital

caliper with a resolution of 0.025 mm. A 10X magnification optical comparator was used

31

to measure the final cross section diameter and the neck radius at the fracture location of

the forged steel specimens. The use of these measurements is further explained in

subsequent sections.

2.1.2.2 Constant amplitude fatigue tests

The uniaxial fatigue tests were performed according to ASTM Standard E606

[2004]. A total of 13 specimens of forged steel and 15 specimens of cast iron were

tested. The standard specifies using a minimum of 10 specimens, so the requirement was

met for both materials. Instron LCF (low cycle fatigue) software was used primarily with

the exception of some of the longer life tests, in which Instron SAX software was used

after switching to load control. During the test the software recorded the total strain

along with the test load at an interval of 2n cycles automatically. Manual data saves were

performed periodically around the expected mid life of the test. A total of seven strain

amplitudes were chosen which included 2%, 1%, 0.5%, 0.35%, 0.25%, 0.2%, and 0.16%.

A minimum of two tests (for each material) were conducted at each strain amplitude with

the exception of 2%, where there was only 1 test for each material. The reason for this

was because the primary interest was the long life region of the curve due to the

application of the material. More specimens were desired for longer life tests.

Strain control was primarily used for the testing with a few exceptions. During

several of the longer life specimens with relatively small strain amplitudes (0.25% or

less) there was a mean stress that built up as the specimen cycled. The tests were

conducted in strain control until the load amplitude became stable and then the test was

conducted in load control at the stable load determined from the strain-controlled test.

For these tests where there was some plastic deformation, the test frequency in load

32

control remained the same as the strain-controlled tests. Load control was also used for

longer life tests where the strain was all or almost all elastic. For these tests the test was

first conducted in strain control to determine if there was any plastic deformation and to

determine the stabilized load. The reason for switching these tests to load control was

due to the frequency limitations of the extensometer. For strain-controlled tests the test

frequencies ranged from 0.1 Hz to 1.0 Hz. For load-controlled tests with little or no

plastic deformation present, the test frequency was increased to 25 Hz to minimize the

time required for each test. A triangular waveform was used in each test.

After each fatigue test, the test specimen was sectioned in the grip section in order

to measure hardness. The specimens were cut using an abrasive cutoff tool with cutting

fluid. The hardness was measured using an Accupro AR-10 Hardness Tester. The

hardness was measured at three locations in the specimen. The hardness values in HRC

were averaged for each specimen. The hardness measured using the Rockwell B and

Rockwell C scales along with the averages for the fatigue specimens are shown in Table

2.2. Both hardness scales were used due to the values measured being at the limits of

both scales. The hardness measurements revealed a hardness value for two forged steel

specimens, listed as FS-2 and FS-11, and one cast iron specimen, CI-12, that were much

lower than the average hardness for the other specimens. The data from these tests also

indicated a higher amount of plastic strain which resulted in a lower stress. This higher

plastic deformation was connected to the lower hardness values. Therefore, these data

points were not used in the determination of any fatigue properties of the forged steel or

cast iron materials. In Table 2.2, the specimens that were not included due to low

hardness values are shown with an asterisk. It should be noted that the forged steel

33

specimen labeled FS-8 also showed a lower hardness value. However the plastic strain

observed during the test was as expected as evidence by the true stress amplitude versus

true plastic strain amplitude plot presented in Section 2.1.3.2.

2.1.3 Experimental results and comparisons

2.1.3.1 Monotonic properties

Properties obtained from the monotonic tensile tests include: modulus of elasticity

(E), yield strength (YS), ultimate strength (Su), percent elongation (%EL), percent

reduction in area (%RA), strength coefficient (K), strain hardening exponent (n), true

fracture strength (σf), and true fracture ductility (εf).

Engineering strain (e) and engineering stress (S) were recorded during the test.

From the engineering stress and strain, the true stress (σ) and true strain (ε), were

calculated using the constant volume assumption which results in the following

relationships:

)1( eS +=σ (2.1)

)1ln( e+=ε (2.2)

True plastic strain (εp) was calculated from:

εσεεεε −=−= ep (2.3)

The Ramberg-Osgood equation is often used to represent the true stress (σ)-true

strain (ε) plot. The Ramberg-Osgood equation is given by:

n

pe KE

1

⎟⎠⎞

⎜⎝⎛+=+= σσεεε (2.4)

34

The strength coefficient (K) and strain hardening exponent (n) are the stress

intercept at a plastic strain of 1 and slope, respectively, to the best fit line of the true

stress (σ) versus true plastic strain (εp) data when plotted in a log-log scale. The equation

of the best fit line is therefore:

( ) npK εσ = (2.5)

The strength coefficient (K) and strain hardening exponent (n) were obtained by

performing a least squares fit of the true stress (σ) versus true plastic strain (εp) data. The

data used in this fit were between the yield stress and the ultimate strength of the

material. This was chosen because of the discontinuous yielding definition in the ASTM

Standard E646 [2004]. The true plastic strain (εp) was the independent variable and the

true stress (σ) was the dependent variable as specified by the ASTM Standard E739

[2004]. The true stress versus true plastic strain plot of both materials is shown in Figure

2.6. It can be seen from the figure that the strength coefficient (K) is slightly higher for

forged steel than for cast iron and the strain hardening exponent is lower for forged steel

than for cast iron.

True fracture strength (σf) can be calculated using the load at fracture, Pf, and the

area at fracture, Af, but when there is necking present there exists a biaxial state of stress

on the cylindrical surface and a triaxial state of stress in the interior of the specimen. In

order to compensate for this state of stress the true fracture strength was calculated using

the Bridgman correction factor which is given by the following equation:

⎟⎠⎞⎜

⎝⎛ +⎟

⎠⎞⎜

⎝⎛ +

=

RD

DR

AP

f

f

f

41ln41 min

min

σ (2.6)

35

where R is the neck radius and Dmin is the minimum diameter of the fracture location.

This was only used for forged steel since the cast iron specimen did not show signs of

necking. For cast iron the true fracture strength (σf) was calculated using the equation:

f

ff A

P=σ (2.7)

where the area at fracture is calculated using the diameter of the specimen after fracture

as measured with an optical comparator.

True fracture ductility (εf) was calculated using the equation:

⎟⎟⎠

⎞⎜⎜⎝

⎛=

ff A

A0lnε (2.8)

where A0 is the initial cross-sectional area.

The monotonic tensile test results for the two materials are summarized in Table

2.3. The monotonic properties for the two materials are shown in Table 2.4. The

monotonic engineering stress-strain curves for the two materials are shown in Figure 2.7.

These stress-strain curves for both materials are shown superimposed on the same plot in

Figure 2.8.

2.1.3.2 Cyclic deformation properties and behavior

The resistance to deformation of a material can change when a cyclic load is

applied rather than a monotonic load. The application of inelastic strain can change the

properties of the material. As a cyclic load is applied, the material may cyclic soften or

cyclic harden. These two terms refer to the decrease and increased resistance to

deformation, respectively. This “cyclic transient behavior” can be observed by plotting

stress amplitude versus the number of cycles. An increase in stress amplitude with

applied strain cycles represents cyclic hardening behavior, while a decrease of stress

36

amplitude represents cyclic softening behavior. The transient response of the ductile cast

iron and the forged steel are shown in semi-log format in Figure 2.9. The response is also

shown normalized in Figure 2.10. One specimen from each strain amplitude is shown in

the plots.

Although this “cyclic transient behavior” does exist, the material stabilizes with

applied cyclic loading. This stabilization is important to the representation of cyclic

material properties. If the material continues to change, material properties would be

dependant on the cycles applied. The midlife of the test was taken as the stabilized state

of the material, except where the test was switched from strain control to load control (for

this case the hysteresis loop at the time of the switch was used). Therefore, the steady

state hysteresis loops from the constant amplitude strain-controlled fatigue tests were

used to determine the following cyclic properties: fatigue strength coefficient (σf’),

fatigue strength exponent (b), fatigue ductility coefficient (εf’), fatigue ductility exponent

(c), cyclic yield strength (YS’), cyclic strength coefficient (K’), and the cyclic strain

hardening exponent (n’). The cyclic properties of ductile cast iron and forged steel are

summarized in Table 2.4.

The steady-state hysteresis loops for the forged steel material and the cast iron

material are shown in Figure 2.11. A summary of the constant amplitude completely

reversed fatigue test data for the forged steel is shown in Table 2.5 and for the cast iron in

Table 2.6.

In order to find the cyclic strength coefficient (K’) and the cyclic strain hardening

exponent (n’), the true plastic strain amplitude (∆εp/2) was calculated using the equation:

Ep

222σεε ∆

−∆

=∆

(2.9)

37

The cyclic strength coefficient (K’) and the cyclic strain hardening exponent (n’) were

obtained by plotting the true stress amplitude (∆σ/2) versus true plastic strain amplitude

(∆εp/2) in log-log scale. The cyclic strength coefficient (K’) is the intercept of stress

amplitude at a plastic strain amplitude of 1, and the cyclic strain hardening exponent (n’)

is the slope of the best fit line. To satisfy the ASTM Standard E739 [2004], the true

plastic strain amplitude (∆εp/2) was the independent variable when performing the least

squares fit of the data. The range of data used to obtain the K’ and n’ values were 0.25%

< ε a< 2% for cast iron and 0.2% < εa < 2% for forged steel. This range represents the

range in which significant plastic deformation occurred. The best fit line of the data is

represented by the equation:

'

2'

2

npK ⎟⎟

⎠

⎞⎜⎜⎝

⎛ ∆=

∆ εσ (2.10)

The K’ and n’ values are used in the Ramberg-Osgood equation that characterizes the

cyclic true stress-strain behavior of the material. The Ramberg-Osgood equation for

cyclic behavior is given by:

'1

'22222

n

KEpe ⎟

⎠⎞

⎜⎝⎛ ∆

+∆

=∆

+∆

=∆ σσεεε (2.11)

The plots of true stress amplitude (∆σ/2) versus true plastic strain amplitude (∆εp/2) in

log-log scale for the two materials along with the best fit lines are shown in Figure 2.12.

Due to the “cyclic transient behavior” the cyclic stress-strain curve is different

than the monotonic curve. The cyclic stress-strain curve was obtained using the applied

strain amplitudes and the stress amplitudes from the stable hysteresis loops. The cyclic

stress-strain curves for the two materials are shown in Figure 2.13. The cyclic stress-

strain curve from both materials are shown superimposed on the same plot in Figure 2.14.

38

The cyclic stress-strain curves are shown superimposed with their respective monotonic

stress-strain curve in Figure 2.15. The cyclic and monotonic stress-strain curves for the

two materials are shown superimposed on the same plot in Figure 2.16. From Figure

2.15(a) it can be seen that the forged steel cyclic softens for the range of available cyclic

stress-strain data. From Figure 2.15(b) it can be seen that the cast iron cyclic hardens.

2.1.3.3 Fatigue behavior and comparisons

When determining strain-life fatigue properties, such as σf’, b, εf’, and c, the

stress amplitude (∆σ/2) and the plastic strain amplitude (∆εp/2) were considered the

independent variables and the fatigue life (2Nf) was considered the dependent variable in

the least squares fit. This is done in accordance with ASTM Standard E739 [2004].

The elastic fatigue behavior of the material can be described by Basquin’s

equation as:

( ) bff N2'

2σσ

=∆ (2.12)

The fatigue strength coefficient (σf’) and the fatigue strength exponent (b) were found by

fitting a line to the true stress amplitude (∆σ/2) versus number of reversals to failure

(2Nf) data in log-log scale. σf’ is intercept at one reversal, 2Nf = 1, and b is the slope of

the best fit line. The range of data used to determine σf’ and b were 0.2% ≤ εa ≤ 2% for

forged steel and 0.16% ≤ εa ≤ 2% for cast iron. The plots of true stress amplitude (∆σ/2)

versus reversals to failure (2Nf) along with the best fit lines for the two materials are

shown in Figure 2.17. Superimposed plots of the two materials are shown in Figure 2.18.

This figure shows that forged steel has a higher fatigue strength than ductile cast iron at

any given life. For a given stress amplitude, the forged steel life is larger by at least an

39

order of magnitude than cast iron at shorter lives, and approximately 50 times larger at

long lives. Since the component is a rotating component in an engine, it is subjected to a

large number of cycles in service. Therefore, the fatigue performance at long life is the

main area of interest. The fatigue limit, considered to be at 106 cycles for both materials,

was 358.9 MPa for forged steel and 262.8 MPa for cast iron. The fatigue strength at 106

cycles for the forged steel material was 36% higher than the fatigue strength of the

ductile cast iron material at the same life.

The relationship between the true plastic strain amplitude and the number of

reversals to failure is given by the Manson-Coffin relationship:

( ) cff

p N2'2

εε

=∆

(2.13)

The fatigue ductility coefficient (εf’) and the fatigue ductility exponent (c) were

determined by fitting a line to the true plastic strain amplitude (∆εp/2) versus reversals to

failure (2Nf) data in log-log scale. εf’ is the intercept at one reversal, 2Nf = 1, and c is the

slope of the best fit line. The range of data used to determine εf’ and c were 0.2% ≤ εa ≤

2% for forged steel and 0.25% ≤ εa ≤ 2% for cast iron. This range was selected as the

range where significant plastic deformation occurred. The plots of true plastic strain

amplitude (∆εp/2) versus reversals to failure (2Nf) along with the best fit lines for the two

materials are shown in Figure 2.19. Superimposed plots of the two materials are shown

in Figure 2.20. Figure 2.20 shows that the forged steel material has a factor of 40 longer

life than the ductile cast iron material for a given plastic strain amplitude in the long life

region.

The total strain is related to the fatigue life by adding the elastic and plastic

portions of the curve. The strain-life equation is given by:

40

( ) ( ) cff

bf

fpea NN

E2'2

'222

εσεε

εε+=

∆+

∆==

∆ (2.14)

The strain-life curves along with the elastic strain portion, plastic strain portion, and

fatigue data for both materials are shown in Figure 2.21. The strain life curve for the two

materials are shown superimposed on the same plot in Figure 2.22. From the figure it can

be seen that the forged steel curve is above the cast iron curve at all lives. In the long life

region, which is the region of importance of this study, there is a factor of approximately

10 between the life of the cast iron and forged steel.

A variation on the strain life curve is the Neuber’s plot. Neuber’s stress range is

calculated by:

( )( ) ( ) ( ) ( ) cbfff

bff NENE ++=∆∆ 2''2'2 22 εσσεσ (2.15)

In Equation 2.15 the term on the left is referred to as Neuber’s parameter. The fatigue

behavior at a notch is often controlled by the stress range and the strain range at the root

of the notch. Neuber’s parameter is significant when comparing the fatigue performance

of crankshaft materials due to the presence of notches, or fillets, in the crankshaft since

this parameter combines the stress range, strain range, and modulus of elasticity. The

Neuber plots for forged steel and cast iron are shown in Figure 2.23. Superimposed

Neuber plots for the two materials are shown in Figure 2.24. From the figure in can be

seen that the forged steel material has superior fatigue performance to the ductile cast

iron material when Neuber’s parameter is used. In the long life region this amounts to a

factor of 50 longer life for the forged steel than the ductile cast iron material for a given

Neuber stress range.

41

2.2 Charpy V-Notch Tests

2.2.1 Specimen and test equipment

The specimen geometry was taken from the ASTM Standard E23 [2004]. The

standard specifies several different geometries which may be used. Of the several

options available, the most commonly used specimen geometry was chosen as the

geometry used for this study. This is the 10mm X 10mm X 55mm specimen geometry

with a v shaped notch which E23 labels as Charpy Impact Test Specimen Type A. The

geometry of the specimens created is shown in Figure 2.25 along with the specified

tolerances.

The specimens used for obtaining impact toughness data by means of the Charpy

impact test were obtained from crankshafts identical to those used to obtain monotonic

and fatigue specimens. The process of forging causes the inclusions to become elongated

in the longitudinal (maximum grain flow) direction of the sample. This elongation of the

inclusions results in lower impact toughness when the notch is oriented in the

longitudinal direction. Therefore, two different specimen orientations were used for the

forged steel specimens. The locations from which the specimens were removed from the

crankshaft are shown in Figure 2.26. Two letter designations are used when referring to

the specimen orientation. The letter “L” represents the longitudinal direction and the

letter “T” represents the transverse direction as indicated in Figure 2.26. One set of

specimens are in the L-T orientation and the other set in the T-L orientation. In this code,

the first letter represents the direction which is normal to the crack plane, and the second

letter designates the direction in which the notch is machined (and the direction of crack

growth). The casting process results in inclusions or porosity which are randomly

42

distributed in the sample and not expected to be elongated in any particular direction.

Therefore, only one set of cast iron specimens were manufactured.

The specimens were machined in The University of Toledo Mechanical,

Industrial, and Manufacturing Engineering Machine Shop. First the rough shape was cut

from the sections shown in Figure 2.26 as discussed previously. Four specimens were

obtained from Section A (as indicated in the figure) of both cast iron and forged steel

crankshafts (L-T). Four specimens were also obtained from each counterbalance section

of the forged steel crankshaft labeled Section B. There are two counterbalance sections

of the crankshaft; therefore 8 specimens in the T-L orientation were obtained from each

crankshaft. The specimens were then machined on a milling machine to an oversized

geometry from that required. The remaining material was removed using a grinding

machine until the specimens were the proper dimensions. The notch was cut using a

horizontal milling machine and a 45 degree double angle milling cutter which had a 0.25

mm radius. All specimen dimensions, including the notch depth and angle were

measured using a 10X magnification optical comparator.

The Charpy impact tests were conducted using a Tinius Olsen pendulum type

impact testing machine. The machine used in the testing is shown in Figure 2.27. The

machine is fitted with a dial indicator that reads directly in energy (kg-m). The machine

was verified prior to being used. A zero verification test was done to ensure that no

uncompensated windage or frictional losses were present in the machine. The test was

conducted with no specimen present, and it was verified that the reading was zero. A

percentage friction and windage loss test was also done as outlined by ASTM Standard

E23 [2004], to verify that the friction and windage loss did not exceed 0.4% of the

43

maximum scale value. The amount of friction and windage loss present in the machine

was within the acceptable range. The pendulum is raised to its initial height and then

released by a lever. The pendulum swings as it impacts the face opposite the notch and

then reaches its final height. The difference between the initial and final heights of the

pendulum results in a reading of absorbed energy on the dial gauge.

2.2.2 Test procedure

Since the impact toughness of a material changes with temperature, an absorbed

energy versus temperature plot is usually constructed. The typical plot has a lower shelf

region and an upper shelf region with a curve connecting the two. Due to the application

of the crankshafts used in this study, less emphasis was put on obtaining the precise upper

and lower shelf regions than conducting the test over a range of temperatures which

would include the operating range of the crankshafts. Six temperature levels were used

for the tests. The temperature levels for all three specimen types (Forged steel L-T, T-L,

and ductile cast iron) ranged between -77°C and 200°C. Two specimens were tested for

each orientation/material at each temperature. Room temperature specimens were tested

first where the room temperature was 26.3°C as measured by a thermometer. For the 0°C

tests, an ice bath was maintained at 0°C. For the tests at -40°C and -77.1°C, a

temperature conditioning bath of dry ice and lab grade isopropanol alcohol was used. For

the tests using a liquid medium the temperature was constantly monitored using a liquid

thermocouple probe and a digital readout. The specimens were immersed in the

temperature conditioning bath for at least 5 minutes prior to testing. The tongs used to

hold the specimen were also immersed in the bath prior to the first test and in between

44

subsequent tests. To ensure even temperature distribution, the conditioning bath was

manually stirred. For the tests conducted at 100°C and 200°C, a Fisher Scientific

Isotemp Oven Model 851F with a digital controller was used. The specimens were

placed in the preheated oven at the specified temperature for at least 1 hour prior to

testing. The tongs were placed in the oven prior to the first test and in between each

subsequent test. For all tests conducted at temperatures other than ambient, the test was

conducted within five seconds of removing the specimen from its temperature

conditioning environment.

2.2.3 Test results and comparisons

The results of the Charpy Impact tests are shown in Table 2.7. The average

absorbed energy values obtained from duplicate Charpy V-notch impact tests at each

temperature are shown as a bar chart in Figure 2.28. From this figure it can be seen that

the forged steel in the L-T direction had a higher absorbed energy value over the entire

range of temperatures. The figure also shows that the ductile cast iron values were the

lowest of the three sets of values over the entire range of temperatures, as expected.

Figure 2.29 shows the absorbed energy versus temperature curve for all three specimen

types. The upper shelf region is shown for the three materials, however with the lowest

temperature at -77°C, the lower shelf region is not clearly shown. The middle transition

curve was obtained by fitting an nth order polynomial to the data. This curve also

indicates that forged steel in the L-T orientation has the highest impact toughness of the

three material/orientations tested regardless of temperature.

45

The fracture surfaces of all specimens tested are shown in Figure 2.30. The

percentage shear fracture (ductile fracture) was also observed for each specimen. The

ASTM Standard E23 Annex 6 [2004] specifies four procedures for determining the

percentage of ductile fracture. The second option, comparing the surface with the

supplied fracture appearance chart, was chosen. The percentage shear, along with the

energy data for each test is given in Table 2.7. The forged steel in both directions

exhibited 100% ductile fracture at 100°C and 200°C. The forged steel specimens also

showed little to no ductile fracture at the sub zero temperatures. The cast iron specimens

exhibited brittle fracture over the entire temperature range tested.

46

Table 2.1: Chemical analysis of the forged steel and ductile cast iron as a percent weight, remaining Fe [Heitmann, 2006].

Element Forged SteelDuctile Cast

Iron C 0.45 3.44

Mn 0.81 0.48 P 0.016 0.019 S 0.024 0.004 Si 0.27 2.38 Al 0.033 0.01 Cr 0.1 0.09 Ni 0.05 0.06 Cu 0.13 0.31 N 0.008 -- O 13 ppm --

47

Table 2.2: Hardness values for (a) forged steel and (b) ductile cast iron monotonic and fatigue specimens.

Specimen FS-2* FS-3 FS-4 FS-6 FS-8 FS-9 FS-10 FS-11* FS-12 FS-13 FS-14 FS-15 18 25 24 27 18 24 23.5 17.5 26 25 20.5 25.5 18 25 24 27 18.5 25 24 18 27 26 22.5 26.5 Measured

HRC 20 25.5 26.5 27.5 18.5 25.5 24.5 19 27.5 27.5 23 27.5

97.5 101.5 100.5 102.5 98 102 101 97 102 103.5 102 102.5 97.5 102 102 102.5 98.5 102 101 97.5 102.5 103.5 102 102.5 Measured

HRB 98.5 103 103 103 99.5 102.5 101 98 103 103.5 102.5 102.5

Average HRC 18.7 25.2 24.8 27.2 18.3 24.8 24.0 18.2 26.8 26.2 22.0 26.5

Average HRB 97.8 102.2 101.8 102.7 98.7 102.2 101.0 97.5 102.5 103.5 102.2 102.5

(a)

Specimen CI-1 CI-2 CI-4 CI-5 CI-6 CI-8 CI-9 CI-10 CI-11 CI-12* CI-13 CI-14 CI-15 CI-16 17 18 16.5 20.5 19.5 19 19 17 19 16 21 18.5 19.5 20

18.5 18 15 19 19 19.5 19.5 18 21 15 19 17 18 18 Measured

HRC 19 18 17 20 21 20 19 18.5 18.5 13.5 20 17 20 20 99 97 94.5 97.5 98 97 98.5 98.5 98.5 94 97.5 98 97 97.5 100 97.5 94.5 97.5 97 96.5 97.5 97 99 94.5 99 97 96.5 98

Measured HRB

100 98 94.5 97 98 97.5 98 96.5 97 93.5 98 97 97 97 Average

HRC 18.2 18.0 16.2 19.8 19.8 19.5 19.2 17.8 19.5 14.8 20.0 17.5 19.2 19.3

Average HRB 99.7 97.5 94.5 97.3 97.7 97.0 98.0 97.3 98.2 94.0 98.2 97.3 96.8 97.5

* Test data were not considered due to low hardness levels

(b)

48

Table 2.3: Result summary of monotonic tensile tests.

Specimen ID

Do, mm (in.)

Df, mm (in.)

Lo, mm (in.)

Lf, mm (in.)

E, GPa (ksi)

YS 0.2% offset, MPa (ksi)

UYS, MPa (ksi)

LYS, MPa (ksi)

YPE, %

Su, MPa (ksi)

K, MPa (ksi)

n %EL %RA R,

mm (in.)

σf* ,

MPa (ksi)

εf

FS-12 5.13 3.33 5.99 9.22 221.3 625.0 681.9 623.8 0.44% 826.6 1,315.6 0.152 54% 58% 1.46 979.5 87% CI-4 5.16 5.00 5.99 6.60 178.2 412.2 -- -- -- 657.6 1,199.0 0.183 10% 6% -- 657.6 6%

* On the forged steel the value of true fracture strength is corrected for necking according to the Bridgman correction factor.

49

Table 2.4: Summary of monotonic and cyclic properties for the two materials.

Monotonic Properties Forged Steel Cast Iron

Ratio

Average Hardness, HRC 23 18

0.8

Average Hardness, HRB 101 97

0.96

Modulus of elasticity, E, Gpa (ksi) 221 (32,088) 178 (25,838)

0.81

Yield Strength (0.2%offset), YS, MPa (ksi) 625 (91) 412 (60)

0.66

Ultimate strength, Su, MPa (ksi) 827 (120) 658 (95)

0.80

Percent elongation, %EL 54% 10%

0.19

Percent reduction in area, %RA 58% 6%

0.10

Strength coefficient, K, MPa (ksi) 1316 (191) 1199 (174)

0.91

Strain hardening exponent, n 0.152 0.183

1.20

True fracture strength, σf, MPa (ksi) 980 (142) 658 (95)

0.67

True fracture ductility, εf 87% 6%

0.07

Cyclic Properties Forged Steel Cast Iron Ratio

Fatigue strength coefficient, σf', MPa (ksi) 1124 163 927 (134) 0.82

Fatigue strength exponent, b -0.079 -0.087 1.10

Fatigue ductility coefficient, εf' 0.671 0.202 0.30

Fatigue ductility exponent, c -0.597 -0.696 1.17

Cyclic yield strength, YS', MPa (ksi) 505 73 519 (75) 1.03

Cyclic strength coefficient, K', MPa (ksi) 1159 168 1061 (154) 0.91

Cyclic strain hardening exponent, n' 0.128 0.114 0.89

Sf = σf'(2Nf)b at Nf = 106, MPa (ksi) 359 (52) 263 (38) 0.73

Average E' Gpa (ksi) 204 (31,437) 174 (25,229) 0.85

Note: Forged steel taken as the base for all ratio calculations

50

Table 2.5: Summary of constant amplitude completely reversed fatigue test results for forged steel.

Spec. ID *

Do, mm (in.)

E',Gpa (ksi)

Testing control mode

Test freq. Hz

∆ε/2, % ∆εp/2

(calc.)%

∆εp/2 (meas.)

%

∆σ/2, MPa (ksi)

σm, MPa (ksi)

N50% , [a] cycles

(Nf)10% , [b]

cycles

(Nf)20% , [c]

cycles

(Nf)50% , [d] cycles

Failure location

[e]

Hardness (HRC)

FS-8 5.18 203.0 strain 0.10 1.981% 1.674% 1.637% 679.7 -3.6 147 263 266 276 IGL 18.3 (0.204) (29,448) (98.6) -(0.5)

FS-14 5.21 191.1 strain 0.50 0.995% 0.711% 0.674% 629.0 -3.7 574 1,001 1,035 1,132 IGL 22 (0.205) (27,709) (91.2) -(0.5)

FS-3 5.18 192.7 strain 0.50 0.999% 0.715% 0.677% 629.3 -2.7 600 1,125 1,142 1,188 IGL 25.2 (0.204) (27,947) (91.3) -(0.4)

FS-9 5.18 207.6 strain 0.83 0.499% 0.253% 0.224% 543.1 14.5 2,450 4,827 4,847 4,894 IGL 24.8 (0.204) (30,115) (78.8) (2.1)

FS-15 5.21 199.1 strain 0.83 0.501% 0.259% 0.232% 534.8 9.3 2,671 4,901 5,056 5,304 IGL 26.5 (0.205) (28,881) (77.6) (1.3)

FS-10 5.13 194.0 strain 0.50 0.349% 0.135% 0.105% 472.7 27.5 8,105 13,515 13,567 13,635 IGL 24 (0.202) (28,131) (68.6) (4.0)

FS-6 5.18 199.7 strain 0.50 0.348% 0.129% 0.124% 485.3 32.1 4,509 8,798 9,127 10,384 IGL 27.2 (0.204) (28,957) (70.4) (4.7)

FS-11* 5.16 207.7 strain 0.83 0.249% 0.080% 0.060% 374.3 37.4 37,345 -- -- 74,691 IGL 18.2 (0.203) (30,128) load (54.3) (5.4)

FS-4 5.21 205.1 strain 0.83 0.251% 0.051% 0.037% 442.7 45.1 55,742 -- -- 111,484 IGL 24.8 (0.205) (29,739) load (64.2) (6.5)

FS-13 5.21 208.5 strain 0.83 0.199% 0.020% 0.008% 396.9 43.0 9,352 -- -- 509,935 IGL 26.2 (0.205) (30,243) load (57.6) (6.2)

FS-2* 5.13 210.5 strain 1.00 0.200% 0.043% 0.035% 346.2 15.8 8,999 -- -- 540,950 IGL 18.7 (0.202) (30,535) load (50.2) (2.3)

FS-16 5.16 NA load 25.00 0.160% 0.000% 0.000% 342.0 0.0 -- -- -- 5,000,000 NA -- (0.203) NA (49.6) (0.0)

FS-7 5.13 228.1 strain 25.00 0.160% 0.000% 0.000% 354.0 1.7 -- -- -- 5,000,000 NA -- (0.202) (33,078) load (51.3) (0.3)

[a] N50% is defined as the midlife cycle.

[b] (Nf)10% is defined as 10% load drop. [c] (Nf)20% is defined as 20% load drop.

[d] (Nf)50% is defined as 50% load drop.

[e] IGL = inside gage length; OGIT = outside gage length but inside test section. * Specimens were not included in fits due to low hardness values

51

Table 2.6: Summary of constant amplitude completely reversed fatigue test results for ductile cast iron.

Speci. ID

Do, mm (in.)

E',Gpa (ksi)

Testing control mode

Test freq., Hz

∆ε/2, % ∆εp/2 (calc.)

%

∆εp/2 (meas.)

%

∆σ/2, MPa (ksi)

σm, MPa (ksi)

N50% , [a]

cycles

(Nf)10% , [b]

cycles

(Nf)20% , [c]

cycles

(Nf)50% , [d] cycles

Failure location

[e]

Hardness, HRC

CI-14 5.21 150.3 strain 0.10 1.994% 1.612% 1.549% 680.1 -25.9 7 NA NA 14 IGL 17.5 (0.205) (21,794) (98.6) -(3.8)

CI-9 5.16 161.8 strain 0.50 1.000% 0.665% 0.634% 595.2 -16.3 35 75 76 76 IGL 19.2 (0.203) (23,459) (86.3) -(2.4)

CI-13 5.18 161.8 strain 0.50 0.995% 0.659% 0.630% 597.9 -20.5 32 89 90 91 IGL 20.0 (0.204) (23,459) (86.7) -(3.0)

CI-15 5.16 177.9 strain 0.50 0.499% 0.212% 0.202% 510.0 -8.2 200 291 313 371 IGL 19.2 (0.203) (25,808) (74.0) -(1.2)

CI-12 5.16 163.7 strain 0.50 0.498% 0.221% 0.201% 492.8 -5.6 450 757 771 789 OGIT 14.8 (0.203) (23,735) (71.5) -(0.8)

CI-8 5.21 179.5 strain 0.50 0.349% 0.083% 0.076% 474.3 10.3 512 975 1,015 1,164 IGL 19.5 (0.205) (26,031) (68.8) (1.5)

CI-11 5.16 174.9 strain 0.83 0.250% 0.021% 0.016% 407.8 12.5 2,916 5,646 5,703 5,770 IGL 19.5 (0.203) (25,368) (59.1) (1.8)

CI-5 5.16 173.9 strain 0.83 0.250% 0.021% 0.022% 407.8 41.3 8,291 -- -- 16,581 IGL 19.8 (0.203) (25,218) load (59.1) (6.0)

CI-10 5.16 176.3 strain 1.00 0.200% 0.008% 0.005% 341.2 30.0 8,184 -- -- 45,105 IGL 17.8 (0.203) (25,575) load (49.5) (4.3)

CI-16 5.21 172.8 strain 1.00 0.199% 0.008% 0.005% 333.5 30.5 7,613 -- -- 57,445 OGIT 19.3 (0.205) (25,065) load (48.4) (4.4)

CI-1 5.21 -- strain 10.00 0.160% 0.000% 0.000% 285.0 0.0 -- -- -- 317,014 IGL 18.2 (0.205) -- load (41.3) (0.0)

CI-6 5.16 -- load 10.00 0.160% 0.000% 0.000% 286.8 0.0 -- -- -- 144,928 IGL 19.8 (0.203) -- (41.6) (0.0)

CI-2 5.18 183.4 strain 0.75 0.160% -

0.001% 0.002% 286.2 20.0 9,377 -- -- 880,814 IGL 18.0 (0.204) (26,598) load 20.00 (41.5) (2.9)

CI-7 5.13 -- load 20.00 0.135% 0.000% 0.000% 240.5 0.0 -- -- -- 5,000,000 NA -- (0.202) -- (34.9) (0.0)

CI-3 5.16 -- load 20.00 0.135% 0.000% 0.000% 240.5 0.0 -- -- -- 5,000,000 NA -- (0.203) -- (34.9) (0.0)

[a] N50% is defined as the midlife cycle.

[b] (Nf)10% is defined as 10% load drop.

[c] (Nf)20% is defined as 20% load drop.

[d] (Nf)50% is defined as 50% load drop.

[e] IGL = inside gage length; AKP = at knife point; OGIT = outside gage length but inside test section. * Specimens were not included in fits due to low hardness values

52

Table 2.7: Summary of results from Charpy impact tests for (a) forged steel L-T, (b) forged steel T-L, and (c) cast iron.

(a) Temperature

(˚C) Absorbed Energy (kg-

m) Joules Percent Shear -77 1.4 13.7 0% -77 1.3 12.7 0% -45 2.5 24.5 10% -45 3.4 33.3 10% 0 4.3 42.2 30% 0 3.8 37.3 40%

26 6.4 62.7 60% 26 5.5 53.9 50% 100 9.5 93.1 100% 100 8 78.4 100% 200 8.8 86.3 100% 200 7.5 73.5 100%

(b) Temperature

(˚C) Absorbed Energy (kg-

m) Joules Percent Shear -77 1.0 9.8 0% -77 1.2 11.8 0% -44 1.3 12.3 0% -44 1.6 15.7 0% 0 2.1 20.6 20% 0 2.4 23.5 10%

26 3.5 34.3 60% 26 4.3 42.2 50% 100 3.2 31.4 100% 100 5.3 52.0 100% 200 5.1 50.0 100% 200 6.6 64.7 100%

(c) Temperature

(˚C) Absorbed Energy (kg-

m) Joules Percent Shear -77 0.3 2.9 0% -77 0.2 2.0 0% -41 0.5 4.9 0% -41 0.3 2.9 0% 0 0.3 2.9 0% 0 0.4 3.9 0%

26 0.5 4.9 0% 26 0.5 4.9 0% 100 0.8 7.8 0% 100 0.9 8.8 0% 200 1.4 13.7 0% 200 1 9.8 0%

53

(a)

(b)

Figure 2.1: Forged steel (a) and ductile cast iron (b) crankshafts used to obtain test specimens.

54

(a)

(b)

Figure 2.2: Photomicrographs of the ductile cast iron material at (a) 500X and (b) 1000X [Laus and Heitmann, 2007].

55

Figure 2.3: Photomicrograph of the forged steel material at 500X.

20 µm

56

Figure 2.4: Specimen geometry for monotonic tensile tests and constant amplitude fatigue tests.

1.125”

3.5

57

(a)

(b)

Figure 2.5: Locations where the monotonic and fatigue specimens were removed from

for forged steel (a) and cast iron (b).

58

100

1000

0.1% 1.0% 10.0%

True Plastic Strain, εp (%)

Tru

e St

ress

, σ (

MPa

) σ=1315.6 (εp)0.1522

K =1315.6MPan = 0.1522R2 = 0.9937

Specimen FS-12

(a)

100

1000

0.1% 1.0% 10.0%

True Plastic Strain, εp (%)

Tru

e St

ress

, σ (

MPa

) σ = 1199(εp)0.1828

K = 1199 MPan = 0.1828R2 = 0.9805

Specimen CI-4

(b)

Figure 2.6: True stress versus true plastic strain for (a) forged steel and (b) ductile cast

iron.

59

0

100

200

300

400

500

600

700

800

900

0% 3% 7% 10%Engineering Strain, e (%)

Eng

inee

ring

Str

ess,

S (M

Pa)

Specimen FS-12

(a)

0

100

200

300

400

500

600

700

0% 3% 7% 10%

Engineering Strain, e (%)

Eng

inee

ring

Str

ess,

S (M

Pa)

Specimen CI-4

(b)

Figure 2.7: Monotonic engineering stress versus strain curves for (a) forged steel and (b)

ductile cast iron.

60

0

100

200

300

400

500

600

700

800

900

0.0 1.0 2.0 3.0 4.0 5.0 6.0 7.0 8.0 9.0 10.0

Engineering Strain (%)

Eng

inee

ring

Str

ess

(MPa

)

CI Monotonic Curve

FS Monotonic Curve

Figure 2.8: Superimposed monotonic engineering stress versus strain curves for forged steel and ductile cast iron.

61

200

300

400

500

600

700

800

1E+0 1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7

Cycles, N

Tru

e St

ress

Am

plitu

de, ∆σ/

2 (M

Pa)

Strain Amplitudes:

(top to bottom)εa = 2.00%εa = 1.00%εa = 0.50%εa = 0.35%εa = 0.25%εa = 0.2%

εa = 0.16%

(a)

200

300

400

500

600

700

800

1E+0 1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7

Cycles, N

Tru

e St

ress

Am

plitu

de, ∆σ/

2 (M

Pa)

Strain Amplitudes:

(top to bottom)εa = 2.00%εa = 1.00%εa = 0.50%εa = 0.35%εa = 0.25%εa = 0.20%εa = 0.16%εa = 0.135%

(b)

Figure 2.9: True stress amplitude versus number of cycles for (a) forged steel and (b)

ductile cast iron.

62

200

300

400

500

600

700

800

0.0 0.2 0.4 0.6 0.8 1.0

Cycle Ratio, (N/Nf)

Tru

e St

ress

Am

plitu

de, ∆σ/

2 (M

Pa)

Strain Amplitudes: (top to bottom)

εa = 2.00%εa = 1.00%εa = 0.50%εa = 0.35%εa = 0.25%εa = 0.2%

εa = 0.16%

(a)

200

300

400

500

600

700

800

0.0 0.2 0.4 0.6 0.8 1.0

Cycle Ratio, (N/Nf)

Tru

e St

ress

Am

plitu

de, ∆

σ/2

(MPa

)

Strain Amplitudes:

(top to bottom)

εa = 2.00%εa = 1.00%εa = 0.50%εa = 0.35%εa = 0.25%εa = 0.20%εa = 0.16%

εa = 0.135%

(b)

Figure 2.10: True stress amplitude versus normalized number of cycles for (a) forged

steel and (b) ductile cast iron.

63

-900

-700

-500

-300

-100

100

300

500

700

900

-3.0% -2.0% -1.0% 0.0% 1.0% 2.0% 3.0%

True Strain, ε (%)

Tru

e St

ress

, σ (

MPa

)

Strain Amplitudes: (starting on

outside)εa = 2.00%εa = 1.00%εa = 0.50%εa = 0.35%εa = 0.25%εa = 0.20%εa = 0.16%

(a)

-900

-700

-500

-300

-100

100

300

500

700

900

-3.0% -2.0% -1.0% 0.0% 1.0% 2.0% 3.0%

True Strain, ε (%)

Tru

e St

ress

, σ (

MPa

)

Strain Amplitudes:

(top to bottom)

εa = 2.00%εa = 1.00%εa = 0.50%εa = 0.35%εa = 0.25%εa = 0.20%εa = 0.16%

(b)

Figure 2.11: Plots of midlife hysteresis loops for (a) forged steel and (b) cast iron.

64

100

1000

0.01% 0.10% 1.00% 10.00%

True Plastic Strain Amplitude, ∆εp/2 (%)

Tru

e St

ress

Am

plitu

de, ∆

σ/2

(MPa

)

Data

Least Squares Fit

∆σ/2 = 1159.4(∆εp/2) 0.1283

K ' = 1159.4 MPan ' = 0.1283R2 = 0.9819

(a)

100

1000

0.01% 0.10% 1.00% 10.00%

True Plastic Strain Amplitude, ∆εp/2 (%)

Tru

e St

ress

Am

plitu

de, ∆

σ/2

(MPa

)

Data

Least Squares Fit

∆σ/2 = 1060.7(∆εp/2) 0.1137

K ' = 1060.7MPan ' = 0.1137R2 = 0.9921

(b)

Figure 2.12: True stress amplitude versus true plastic strain amplitude for (a) forged steel

and (b) ductile cast iron.

65

0

100

200

300

400

500

600

700

800

0.0% 0.5% 1.0% 1.5% 2.0% 2.5%True Strain Amplitude, ∆ε/2 (%)

Tru

e St

ress

Am

plitu

de, ∆

σ/2

(MPa

)

Data

Cyclic Stress-Strain Equation

(a)

0

100

200

300

400

500

600

700

800

0.0% 0.5% 1.0% 1.5% 2.0% 2.5%True Strain Amplitude, ∆ε/2 (%)

Tru

e St

ress

Am

plitu

de, ∆

σ/2

(MPa

)

Data

Cyclic Stress-Strain Equation

(b)

Figure 2.13: True stress amplitude versus true strain amplitude for (a) forged steel and

(b) ductile cast iron.

66

0

100

200

300

400

500

600

700

800

0.0% 0.5% 1.0% 1.5% 2.0% 2.5%

True Strain Amplitude ∆ε/2 (%)

Tru

e St

ress

Am

plitu

de ∆σ/

2 (M

Pa)

Forged Steel

Cast Iron

Figure 2.14: Superimposed cyclic stress-strain curves for forged steel and ductile cast iron.

67

0

100

200

300

400

500

600

700

800

0.0% 1.0% 2.0% 3.0%True Strain (%)

Tru

e St

ress

(M

Pa)

Cyclic C

Monotonic C

(a)

0

100

200

300

400

500

600

700

800

0.0% 1.0% 2.0% 3.0%True Strain (%)

Tru

e St

ress

(M

Pa)

Cyclic Curve

Monotonic Curve

(b)

Figure 2.15: Superimposed plots of monotonic and cyclic true stress versus true strain

curves for (a) forged steel and (b) ductile cast iron.

68

0

100

200

300

400

500

600

700

800

900

0.0 1.0 2.0 3.0 4.0 5.0 6.0 7.0 8.0 9.0 10.0

True Strain (%)

Tru

e St

ress

(M

Pa)

Forged Steel Monotonic

Forged Steel Cyclic

Cast Iron Cyclic

Cast Iron Monotonic

Figure 2.16: Superimposed plots of monotonic and cyclic true stress versus true strain curves for forged steel and ductile cast iron.

69

100

1000

1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure, 2Nf

Tru

e St

ress

Am

plitu

de, ∆

σ/2

(MPa

)

Fatigue Data

Least Suare Fit

∆σ/2 =1124.3 (2Nf)-0.0787

σf ' = 1124.3 MPab = -0.0787R2 = 0.9525

(a)

100

1000

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure, 2Nf

Tru

e St

ress

Am

plitu

de, ∆

σ/2

(MPa

)

Fatigue Data

Least Suare Fit

∆σ/2 =926.8(2Nf) -0.0869

σf ' = 926.8 MPab = -0.0869R2 = 0.9810

(2)

(b)

Figure 2.17: True stress amplitude versus reversals to failure for (a) forged steel and (b)

ductile cast iron.

70

100

1000

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure, 2Nf

Tru

e St

ress

Am

plitu

de (M

Pa)

(2)

Forged Steel

Cast Iron

Figure 2.18: Superimposed plots of true stress amplitude versus reversals to failure for forged steel and ductile cast iron.

71

0.01%

0.10%

1.00%

10.00%

1E+2 1E+3 1E+4 1E+5 1E+6 1E+7

Reversals to Failure, 2Nf

Tru

e Pl

astic

Str

ain

Am

plitu

de, ∆

ε p/2

(%)

Fatigue Data

Least Squares Fit

∆εp/2 = 0.6707(2Nf) -0.5971

εf ' = 0.6707c = -0.5971R2 = 0.9879

(a)

0.01%

0.10%

1.00%

10.00%

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7

Reversals to Failure, 2Nf

Tru

e Pl

astic

Str

ain

Am

plitu

de, ∆

ε p/2

(%)

Fatigue DataLeast Squares Fit

∆εp/2 = 0.2023(2Nf) -0.6959

εf ' = 0.2023c = -0.6959R2 = 0.9803

(b)

Figure 2.19: True plastic strain amplitude versus reversals to failure for (a) forged steel

and (b) ductile cast iron.

72

0.01%

0.10%

1.00%

10.00%

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7

Reversals to Failure, 2Nf

Tru

e Pl

astic

Str

ain

Am

plitu

de, ∆

ε p/2

(%)

Forged Steel

Cast Iron

Figure 2.20: Superimposed plots of true plastic strain versus reversals to failure for forged steel and ductile cast iron.

73

0.01%

0.10%

1.00%

10.00%

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure, 2Nf

Tru

e St

rain

Am

plitu

de, ∆

ε/2,

%

Strain-Life Equation

Elastic Strain

Plastic Strain

Fatigue Data (Plastic)

Fatigue Data (Elastic)

Fatigue Data (Total)

∆ε / 2

∆ε p /2

∆ε e / 2 (2)

(a)

0.01%

0.10%

1.00%

10.00%

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure, 2Nf

Tru

e St

rain

Am

plitu

de, ∆

ε/2,

%

Strain-Life Equation

Elastic Strain

Plastic Strain

Fatigue Data (Plastic)

Fatigue Data (Elastic)

Fatigue Data (Total)

∆ε / 2

∆ε p /2

∆ε e / 2 (2)

(b)

Figure 2.21: True strain amplitude versus reversals to failure for (a) forged steel and (b)

ductile cast

74

0.01%

0.10%

1.00%

10.00%

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure, 2Nf

Tru

e St

rain

Am

plitu

de, ∆

ε/2,

%

(2)

(2)

Forged Steel

Cast Iron

Figure 2.22: True strain amplitude versus reversals to failure for forged steel and ductile cast iron.

75

100

1000

10000

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure, 2Nf

Neu

ber

Stre

ss R

ange

[(∆ε

)(∆σ

)E]1/

2 , Mpa

Neuber Data

Fitted Equation

(2)

(a)

100

1000

10000

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure , 2Nf

Neu

ber

Stre

ss R

ange

[(∆ε

)(∆σ

)E]1/

2 , M

Pa

(2)

(b)

Figure 2.23: Neuber stress range versus reversals to failure for (a) forged steel and (b)

ductile cast iron.

76

100

1000

10000

1E+1 1E+2 1E+3 1E+4 1E+5 1E+6 1E+7 1E+8

Reversals to Failure, 2Nf

Neu

ber

Stre

ss R

ange

[(∆ε

)(∆σ

)E]1/

2 , MPa

(2)

(2)

Forged Steel

Cast Iron

Figure 2.24: Superimposed Neuber stress range versus reversals to failure for forged steel and ductile cast iron.

77

Figure 2.25: Charpy impact specimen geometry.

Figure 5.XX Locations of the Charpy impact specimen orientations

Longitudinal direction

Transverse direction

Section B T-L Specimens Section A

L-T Specimens

Figure 2.26: Locations on the crankshaft where Charpy v-notch specimens were machined from.

78

Figure 2.27: Tinius Olsen impact test machine.

79

Figure 2.28: Average absorbed energy values at the different test temperatures for forged

steel (L-T, T-L) and ductile cast iron.

0

10

20

30

40

50

60

70

80

90

100

-100 -50 0 50 100 150 200 250

Test Temperature (˚C)

Abs

orbe

d E

nerg

y (J

)

Forged steel L-T

Forged steel T-L

Cast Iron

Figure 2.29: Absorbed energy versus test temperature for forged steel (L-T, T-L) and

cast iron specimens.

80

(a)

(b)

(c)

Figure 2.30: Fracture surfaces of (a) forged steel L-T, (b) forged steel T-L, and (c) ductile cast iron specimens in order of increasing temperature from left to right.

81

CHAPTER 3

COMPONENT TESTING PROCEDURES AND RESULTS

3.1 Test Apparatus and Procedures

3.1.1 Loading conditions and test fixture

In order to compare the fatigue lives of the forged steel and ductile cast iron

crankshafts, constant amplitude, load-controlled fatigue tests were performed on the two

crankshafts. A crankshaft experiences two types of in service loading, bending and

torsion. Previous studies and the dynamic load analysis of the engine showed that the

effect of torsion was negligible compared to bending. Since bending was the primary in

service loading mechanism, it was used as the loading mechanism for the component

fatigue tests. In order to minimize the magnitude of the applied loads necessary to

achieve the desired stress levels, cantilever bending was used, rather than three-point

bending. While the cantilever bending fixture did minimize the loads required, it did

decrease the rigidity of the test fixture. This increased deflection was the limiting factor

in the test frequency.

Finite element analysis in conjunction with the dynamic loading analysis of the

engine identified the critical location, or highest stresses location, of the crankshafts. For

details on identification of the critical location refer to Chapter 4. In order to isolate the

critical location in the crankshafts, the crankshaft was tested such that the critical

82

crankpin fillet had a higher bending moment than the non-critical crankpin fillet. This

was accomplished by orienting the crankshaft such that the critical crankpin fillet had a

longer moment arm. Figure 3.1 shows the production, as tested forged steel crankshaft

with the critical location identified. Figure 3.2 shows the same for the ductile cast iron

crankshaft.

The test fixture was designed so there was a vertical support that clamped onto the

main bearing section of the crankshaft and the load was applied through a moment arm

attached to the nose of the crankshaft. The test fixture was machined from a solid bar of

4” by 3” steel. The vertical support of the test fixture was welded to a ¾” thick plate of

steel so that the fixture could be bolted to the machine test bed. A hole was bored into

the vertical support that had the same diameter as the main bearing sections of both

crankshafts. Identical diameters allowed the test fixture to be used for both crankshafts.

The moment arm was machined from the same 4” by 3” bar steel. A hole the precise size

of the nose section of the crankshafts was bored into the end of the moment arm for

attachment to the crankshaft. Clamping the crankshaft was accomplished with four ½”

diameter bolts on both the vertical support and the moment arm. All bolts were tightened

with a dial gage torque wrench to the same torque prior to testing to ensure an evenly

distributed clamping force. A schematic of the test fixture and set-up is shown in Figure

3.3, where the forged steel crankshaft is shown only as an example. The test set-up for

the forged steel crankshaft is shown in Figure 3.4 and for the cast iron crankshaft in

Figure 3.5.

A rod end bearing was used to apply the load to minimize any misalignment in

the test set-up. As mentioned previously, bending was the only loading desired for the

83

test. However, often times in applying a bending load, an axial force is also present. In

order to minimize this unwanted axial force, the motion between the moment arm and the

loading applicator was not constrained horizontally. Slots were machined into the

loading end of the moment arm. A rod fitted with needle roller bearings was attached to

the rod end bearing. The roller bearings were necessary to minimize the frictional force

that would be present if the rod was allowed to slide in the moment arm slots, rather than

roll. Figure 3.6 shows the rod end bearing and slotted end of the moment arm with the

roller bearings. To verify that there was no significant axial force present, a crankshaft

was fitted with strain gages and loads were applied. The results were compared to those

obtained from FEA and analytical (i.e. Mc/I) calculations. The results showed that there

was no axial force present. The verification of the test fixture is presented in Chapter 4 as

part of the stress analysis.

3.1.2 Test procedures

The bending fatigue tests were performed at room temperature, which was

monitored with a digital thermometer and recorded for each test along with the humidity

which was measured using a precision hydrometer. Tests were conducted using a

sinusoidal waveform and constant amplitude load control. Test frequencies between 1.4

and 3 Hz were used for all tests, with the lower frequency used for the higher load levels

and the higher frequency used for lower load levels. The stress ratio, or R-ratio, is the

ratio of maximum stress to minimum stress. The dynamic load analysis that was

performed resulted in load versus crank angle data for both crankshafts as shown by

Montazersadgh and Fatemi [2007]. As the crankshaft rotates through the engine cycles,

84

the loading, which is primarily bending, changes. The loads are also dependent on the

rpm of the engine. The dynamic analysis showed that the average ratio of minimum load

to maximum load was approximately equal to -0.2 for both crankshafts [Montazersadgh,

2007]. Therefore, an applied R ratio of -0.2 was used for all tests.

Four load levels were used for both crankshafts, with three tests at each load level

to assess variability and scatter. The load levels resulted in lives between 103 and 4 x 106

cycles. The parameters for each test, along with the results are summarized in Table 3.1.

In the table, a positive load value is applied upwards, causing compressive stress on the

top of the crankshaft and tensile stress at the bottom where the critical section was

located.

The forged steel and ductile iron crankshafts were designed to operate in very

similar engines. The crankshafts were of similar size and similar mass, 3.9 kg for the

forged steel and 3.7 kg for the cast iron crankshaft. Due to their similar size and

application, both crankshafts would experience similar in service loading. Therefore,

both crankshafts were tested at the same bending moment levels.

3.2 Failure Criterion

Initially the crankshafts were tested until the point which they could no longer

maintain the applied cyclic load. At this point the displacement versus cycles curve

reached an asymptotic value and the crankshaft was considered to be fractured. After

several tests were completed, it was found that the crack growth life of the component

was one-half to two-thirds the life of the component. Failure of the crankshaft could,

therefore, not be the point when the crankshaft fractured. The crankshaft, being a

85

rotating engine component, could not function in a state of increased deflection. Crack

initiation was used as a failure criterion for the crankshafts. In order to use crack

initiation as a failure criterion, however, the point at which the crack initiated must be

clearly identifiable. Each test was stopped at intervals corresponding to ten to twenty

percent of the expected life of the component, at which point the crankshaft was

inspected for the presence of a crack. If no crack was present the test was continued. If a

crack was present, the crack was measured and recorded. A light white coating of paint

was applied to the critical fillet area in order to help detect a crack, as shown in Figure

3.7. The physical crack length was monitored using putty that was molded to the cracked

area and then removed leaving a raised imprint of the crack, as shown in Figure 3.8. The

crack length was then measured from the putty using a digital caliper. As the crack

grew, the test was stopped at intervals corresponding to ten to twenty percent of the

expected life and the crack length was measured and recorded. It was found that by the

time the crack was detected it was on the order of 6 mm or longer. Crack initiation, for

life prediction purposes, is usually defined as a crack on the order of 1 mm or 2 mm.

Therefore, the point at which a crack was visually detected was not a desired definition of

initiation point, since the crack was already long at this point.

When a crack was present and as it grew, there was, as expected, a decrease in

stiffness, or in other words an increase in deflection. Using the measured crack length

data, along with the recorded displacement amplitude data for each test, a correlation

between the change in displacement amplitude and the crack length was developed for

both forged steel and ductile cast iron crankshafts. For both materials, the change in

displacement amplitude versus measured crack length was plotted. The base for

86

determining the change in displacement amplitude for each test was taken as the point

where the displacement amplitude was stable, as indicated by a horizontal line in Figure

3.9.

The change in displacement amplitude versus measured crack length plot is

shown in Figure 3.10 for the forged steel crankshaft and in Figure 3.11 for the cast iron

crankshaft. In Figures 3.10 and 3.11 each test is shown in a different color. The data are

shown superimposed in Figure 3.12.

Measured crack length versus cycle data for the forged steel crankshafts are

shown in Figure 3.13 and for the cast iron crankshafts in Figure 3.14. For a given

change in displacement amplitude the size of a crack can be determined from Figures

3.10 and 3.11. Then from Figures 3.13 and 3.14, knowing the size of the crack, the

number of cycles can be determined. This method allows for the change in displacement

amplitude versus crack length data to be extrapolated to find the change in displacement

amplitude corresponding to a given length. Then the cycles where the crack initiation

(i.e. a crack on the order of 2 mm) occurred can be determined from the data.

Using the fitted equation for each crankshaft, the change in displacement

amplitude was extrapolated for a crack length of 2 mm. From the relationship between

the change in displacement amplitude and crack length, a crack length of 2 mm would

result in a very small change in displacement amplitude. The change in displacement

amplitude for a 2 mm crack was on the order of a micrometer, such that the position

transducer of the test frame could not accurately detect this change. However, when

there was any recorded increase in displacement amplitude the relationship suggests that

there was a crack present. The data from each test was analyzed and the cycle at which

87

there was a measurable increase in displacement amplitude was determined to be the

crack initiation point. An expanded scale plot of the displacement amplitude versus

cycles plot for the forged steel crankshafts is shown in Figure 3.15 and in Figure 3.16 for

the ductile cast iron crankshafts. These figures show that the displacement amplitude is

relatively steady for a period in the test after the full load is applied and prior to the

formation of a crack. At the point where a crack develops, an increase in the

displacement amplitude is observed.

The predicted crack lengths obtained from the change in displacement amplitude

versus crack length plots shown in Figures 3.10 and 3.11 were compared with the

measured crack lengths. The predicted crack length versus measured crack length is

shown in Figure 3.17 for the forged steel crankshafts and in Figure 3.18 for the ductile

cast iron crankshafts. Scatter bands at factors of plus and minus two are also plotted.

The figures show that the predicted crack lengths are within a factor of two of the

measured crack lengths for both the forged steel and ductile cast iron crankshafts.

A change in displacement amplitude of 5% was also used as a failure criterion for

comparison purposes. A 5% change was much more apparent than the small change in

displacement amplitude that was used for determining crack initiation. From the

displacement amplitude versus cycles plot shown in Figure 3.9, it can be seen that the

displacement amplitude reached a constant value while no crack was present and then

began to increase as the crack grew. The figures also show that for some of the short life

tests the displacement amplitude curves reached an asymptotic value; this was

determined to be the point of fracture. This fracture point was only reached for several

tests due to the significant amount of time it took to grow the crack to that length.

88

3.3 Results and Comparisons

The moment amplitude versus cycles to failure for both forged steel and cast iron

crankshafts using the crack initiation criterion is shown in Figure 3.19. The figure shows

that for a given applied moment, the life of the forged steel crankshaft is approximately

six times longer than the life of the ductile cast iron crankshaft. The moment amplitude

versus cycles to failure for the forged steel and ductile cast iron crankshafts using the 5%

change in displacement amplitude criterion is shown in Figure 3.20. Using the 5%

change in displacement criterion, the difference in life is less at shorter lives when

compared to the crack initiation criterion, but the difference is greater at long lives. At

long life, there is approximately an order of magnitude difference between the life of the

forged steel and cast iron crankshafts. The divergence of the curves at longer lives

suggests that the forged steel crankshaft had a slower crack growth rate than the ductile

cast iron crankshaft.

A fatigue limit is important for a long life component; if the loads or stresses are

below the fatigue limit, failure will likely not occur. When tested at a moment amplitude

of 431 N-m, the forged steel crankshaft had a life greater than 4 x 106 cycles and was

considered a run-out. Two forged steel crankshafts were tested at this level with no

failures. The cast iron crankshafts at this load level failed between 75,200 and 82,200

cycles. The issue of fatigue limit is important when comparing the fatigue lives for a

long life component such as a crankshaft. The cast iron crankshaft has fatigue strength of

316 N-m at 106 cycles based on either the crack initiation or 5% increase in displacement

amplitude criteria. This suggests a 36% higher fatigue strength for the forged steel

89

crankshaft as compared to the cast iron crankshaft. The fatigue strength at 106 cycles for

the two materials obtained from specimen tests in Chapter 2 show the fatigue strength of

the cast iron to be 263 MPa and for the forged steel to be 359 MPa. This also suggests a

36% higher fatigue strength for the forged steel as compared to the cast iron. Therefore,

the differences in fatigue strengths of the cast iron and forged steel for the components

and for the materials are the same. One contributing factor to the same ratio is that the

geometries of the two crankshafts were very similar, therefore the primary difference in

the two crankshafts for component testing purposes was the difference in material

properties.

The literature suggests that fatigue behavior of cast iron is similar to steel and

therefore cast iron should also have a fatigue limit at about 106 cycles [Juvinall and

Marshek, 1991]. The existence of a fatigue limit for the cast iron crankshaft could not be

verified with the limited number of components available for testing and the length of

time required for high cycle testing.

As shown in Figures 3.19 and 3.20, the scatter in the component fatigue life test

data is small for both the crack initiation and 5% change in displacement criteria (within

a factor of about 2 for the forged steel crankshaft and a factor of about 3 for the cast iron

crankshaft). The scatter was similar to what was seen in the specimen fatigue tests. For

both the specimen tests and the component tests, the scatter for the cast iron was

somewhat better than expected. Porosity that is typically present in castings can

contribute to increased scatter due to its random distribution and size. Cracks can

develop from the porosity and therefore their randomness in size and distribution can

influence the fatigue life.

90

Based on the specimen tests conducted on the cast iron specimens, the cast iron

exhibits a cyclically hardening behavior as shown in Chapter 2. The component tests on

the cast iron crankshafts also showed a cyclically hardening behavior. The hardening

behavior does not, however, directly correlate to the hardening behavior observed in the

specimen tests. In the specimen tests the entire gage section experiences the cyclically

hardening. In the component tests the stresses in the crankshaft are completely elastic

with the exception of the fillets where there is plastic deformation. Therefore, the entire

crankshaft is not cyclically hardening, but rather there is local hardening in the highly

stressed fillet locations. Since this hardening is localized, the amount of hardening is not

as large as would be seen in specimen testing. This hardening behavior is shown in

Figures 3.21 and 3.22. The figure shows that the displacement amplitude for the cast iron

crankshafts decreases at the beginning of the test before reaching a stable value. The

forged steel crankshafts showed neither a cyclically hardening nor softening behavior, as

expected from the specimen tests.

The ultimate cause of failure for each crankshaft was a crack that developed and

grew in the critical location (crankpin fillet). Therefore all crankshafts that failed, had the

same failure location. On one forged steel and several cast iron crankshafts a secondary

crank grew in the crankpin fillet opposite to the critical crankpin fillet. These cracks

developed after the cracks at the critical location and were not considered the cause of

failure. The crankshafts where these secondary cracks developed are indicated in Table

3.1.

A typical fatigue fracture of the forged steel crankshaft is shown in Figures 3.23

and Figure 3.24. A typical fatigue fracture for the cast iron crankshaft is shown in

91

Figures 3.25 and 3.26. The figures show that the fracture surface is smoother for the

forged steel when compared to the cast iron. The figures also show that the crack grew

through the circular cross section for the cast iron crankshaft, but not for the forged steel

crankshaft. The crack in the forged steel crankshaft grew approximately half way

through the circular cross section and then veered off at approximately a 45 degree angle

which is attributed to the geometry of the component. The component was no longer the

weakest through the circular cross section once the crack grew long.

The cross section picture for the cast iron crankshaft shows the eccentricity of the

center hole in the crankpin. This eccentricity increased the amount of material in the

highest stressed location of section, at the bottom of Figure 3.25. The eccentricity of the

oil bore was accounted for in the finite element model and also the analytical stress

calculations presented in Chapter 4.

92

Table 3.1: Test parameters and results for the forged steel and ductile cast iron crankshaft fatigue tests.

Forged Steel Crankshaft

Crank ID

Freq (Hz)

R-Ratio

Ma (N-m)

Pmax (kN)

Pmin (kN) Pa (kN) Pm

(kN)

Observed Crack Length

(mm)

N, Crack Observed

Nf, Crack Initiation from Disp. Data

Extrapolation

Nf, 5% Change in

Disp. Amp.

Failure Location

FS-2 1.4 -0.2 630 2.67 -0.53 1.60 -1.07 44.96 98,198 29,248 45,568 1 FS-3 1.4 -0.2 630 2.67 -0.53 1.60 -1.07 51.82 120,492 45,302 69,670 1 FS-4 1.4 -0.2 630 2.67 -0.53 1.60 -1.07 -- -- 58,236 90,853 1 FS-5 2.5 -0.2 517 2.18 -0.44 1.31 -0.87 10.52 165,000 145,000 234,289 1 FS-6 2.5 -0.2 517 2.18 -0.44 1.31 -0.87 11.71 120,000 98,741 213,885 1 FS-7 2.5 -0.2 517 2.18 -0.44 1.31 -0.87 -- -- 204,174 396,011 1 & 2 FS-9 3 -0.2 431 1.82 -0.36 1.09 -0.73 None >2,090,000 Runout No crack

FS-10 3 -0.2 431 1.82 -0.36 1.09 -0.73 None >3,980,000 Runout No crack FS-8 3 -0.2 350 1.48 -0.30 0.89 -0.59 None >3,240,000 Runout No crack

Cast Iron Crankshaft CI-2 1.4 -0.2 630 2.67 -0.53 1.60 -1.07 11.43 11,504 7,132 17,353 1 CI-3 1.4 -0.2 630 2.67 -0.53 1.60 -1.07 5.08 11,692 9,256 17,380 1 CI-4 1.4 -0.2 630 2.67 -0.53 1.60 -1.07 3.175 8,021 8,021 20,957 1 CI-5 2 -0.2 517 2.18 -0.44 1.31 -0.87 8.51 31,464 25,512 47,513 1 & 2 CI-6 2 -0.2 517 2.18 -0.44 1.31 -0.87 8.81 34,898 24,096 52,790 1 & 2 CI-7 2 -0.2 517 2.18 -0.44 1.31 -0.87 12.55 42,750 37,380 54,966 1 & 2 CI-1 2.5 -0.2 431 1.82 -0.36 1.09 -0.73 7.62 113,043 75,200 132,877 1 CI-9 2.5 -0.2 431 1.82 -0.36 1.09 -0.73 13.60 90,175 78,367 121,866 1

CI-10 2.5 -0.2 431 1.82 -0.36 1.09 -0.73 37.06 -- 82,200 143,259 1 & 2 CI-8 2.5 -0.2 350 1.48 -0.30 0.89 -0.59 19.79 985,496 920,783 1,005,665 1 & 2

CI-11 2.5 -0.2 350 1.48 -0.3 0.8896 -0.59 32.72 -- 301,774 370,216 1 & 2

2

1

93

Figure 3.1: Forged steel crankshaft in its final machined condition.

Figure 3.2: Ductile cast iron crankshaft in its final machined condition.

Critical fillet

Critical fillet

94

Figure 3.3: Schematic of test set-up.

Figure 3.4: Test set-up for the forged steel crankshaft.

95

Figure 3.5: Test set-up for the ductile cast iron crankshaft.

Figure 3.6: Close up of load application area of moment arm showing rod end bearing and roller bearings.

Needle roller bearings

Rod end bearing

96

Figure 3.7: Critical fillet area of crankshaft painted to better observe crack.

Figure 3.8: Imprint of crack with putty.

Crack Length

97

1

2

3

4

5

6

7

8

0 200,000 400,000 600,000 800,000 1,000,000

Cycles (N)

Dis

plac

emen

t Am

plitu

de (m

m)

630 N-m

517 N-m 431 N-m 350 N-m

(a)

1.5

2

2.5

3

3.5

4

4.5

5

5.5

6

0 200000 400000 600000 800000 1000000

Cycles (N)

Dis

plac

emen

t Am

plitu

de (m

m)

630 N-m

431 N-m

517 N-m

350 N-m

(b)

Figure 3.9: Displacement amplitude versus number of cycles for the (a) forged steel

crankshafts and (b) ductile cast iron crankshafts.

98

0

1

2

3

4

5

6

7

0 10 20 30 40 50 60 70 80Crack Length (mm)

Cha

nge

in D

ispl

acem

ent A

mp.

(mm

) FS-2 FS-3

FS-5 FS-6

630 N-m

517 N-m

y = 2E-06 x 3.4382

Figure 3.10: Change in displacement amplitude versus crack length for the forged steel crankshafts.

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

2

0 10 20 30 40 50 60 70 80

Crack Length (mm)

Cha

nge

in D

ispl

acem

ent A

mp.

(mm

)

CI-2 CI-3 CI-4CI-5 CI-6 CI-7CI-1 CI-9 CI-10CI-8 CI-11

630 N-m517 N-m431 N-m350 N-m

y = 0.0004 x 1.9

Figure 3.11: Change in displacement amplitude versus crack length for the cast iron

crankshafts.

99

0

1

2

3

4

5

6

7

0 10 20 30 40 50 60 70 80Crack Length (mm)

Cha

nge

in D

ispl

acem

ent A

mp.

(mm

)

FS-2 FS-3FS-5 FS-6CI-1 CI-2CI-3 CI-4CI-5 CI-6CI-7 CI-8CI-9 CI-10CI-11

Figure 3.12: Superimposed plot of change in displacement amplitude versus crack length for the forged steel and cast iron crankshafts.

100

0

10

20

30

40

50

60

70

0 100,000 200,000 300,000 400,000

Cycles

Cra

ck L

engt

h (m

m)

FS-2 FS-3

FS-5 FS-6

Figure 3.13: Measured crack length versus cycles for the forged steel crankshafts.

0

10

20

30

40

50

60

70

80

0 200,000 400,000 600,000 800,000 1,000,000Cycles

Cra

ck L

engt

h (m

m)

CI-2 CI-3 CI-4

CI-5 CI-6 CI-7

CI-1 CI-9 CI-10

CI-8 CI-11

Figure 3.14: Measured crack length versus cycles for the ductile cast iron crankshafts.

101

2.4

2.45

2.5

2.55

2.6

2.65

2.7

2.75

0 50000 100000 150000 200000

Cycles (N)

Dis

plac

emen

t Am

plitu

de (m

m)

Crack initiation

Figure 3.15: Displacement amplitude versus cycles for a forged steel crankshaft with the

crack initiation point highlighted.

2.75

2.8

2.85

2.9

2.95

3

3.05

3.1

3.15

3.2

3.25

0 50000 100000 150000 200000

Cycles (N)

Dis

plac

emen

t Am

plitu

de (m

m)

Crack initiation

Figure 3.16: Displacement amplitude versus cycles for a ductile cast iron crankshaft with

the crack initiation point highlighted.

102

0

10

20

30

40

50

60

70

80

0 20 40 60 80

Measured Crack Length (mm)

Pred

icte

d C

rack

Len

gth

(mm

)

Figure 3.17: Predicted crack length versus measured crack length for the forged steel

crankshafts. The same symbols correspond to crack lengths of the same crankshaft.

0

10

20

30

40

50

60

70

80

0 20 40 60 80

Measured Crack Length (mm)

Pred

icte

d C

rack

Len

gth

(mm

)

Figure 3.18: Predicted crack length versus measured crack length for the ductile cast iron

crankshafts. The same symbols correspond to crack lengths of the same crankshaft.

103

y = 2555.8x-0.1331

R2 = 0.8128

y = 2147.3x-0.139

R2 = 0.9535

100

1000

1.E+03 1.E+04 1.E+05 1.E+06 1.E+07

Cycles to Failure (Nf)

Mom

ent A

mpl

itude

(N-m

)

Forged Steel

Cast Iron

Figure 3.19: Moment amplitude versus cycles to failure using the crack initiation failure criterion.

y = 2401.8x-0.1218

R2 = 0.8656

y = 3115.8x-0.1652

R2 = 0.9579

100

1000

1.E+03 1.E+04 1.E+05 1.E+06 1.E+07

Cycles to Failure (Nf)

Mom

ent A

mpl

itude

(N-m

)

Forged Steel

Cast Iron

(3)

Figure 3.20: Moment amplitude versus cycles to failure using the 5% change in displacement amplitude failure criterion.

104

2

2.5

3

3.5

4

4.5

0 1000 2000 3000 4000 5000 6000 7000 8000

Cycles (N)

Dis

plac

emen

t Am

plitu

de (m

m)

CI 630 N-m

CI 431 N-m

CI 517 N-m

CI 350 N-m

Figure 3.21: Cast Iron displacement amplitude versus cycles plot showing hardening behavior.

2.84

2.86

2.88

2.9

2.92

2.94

0 1000 2000 3000 4000 5000 6000

Cycles (N)

Dis

plac

emen

t Am

plitu

de (m

m)

Figure 3.22: Expanded view of the displacement amplitude versus cycles plot for a cast iron crankshaft tested at 431 N-m.

105

Figure 3.23: Example of a typical fatigue fracture surface for the forged steel crankshaft.

Figure 3.24: Side view of typical fatigue fractured forged steel crankshaft.

Crack initiation site

Crack growth direction

106

Figure 3.25: Example of a typical fatigue fracture surface for the cast iron crankshaft.

Figure 3.26: Side view of typical fatigue fractured cast iron crankshaft.

Crack initiation site

Crack growth direction

107

CHAPTER 4

STRESS ANALYSIS AND FATIGUE LIFE PREDICTIONS

4.1 Analytical Stress Calculations

To verify the nominal stress obtained from FEA which was used for the life

predictions, analytical stress calculations were performed. Since the component was

loading in bending only, the stress was calculated using the equation:

IcMS a

a = (4.1)

where Sa is the stress amplitude, Ma is the moment amplitude, c is the distance from the

centroid to the location where the stress is calculated and I is the area moment of inertia.

The forged steel crankshaft had an outer crankpin diameter of 3.68 cm and an inner

diameter of 1.70 cm, which were used to calculate the area moment of inertia, which is

8.59 cm4.

For the cast iron crankshaft, the calculation was slightly more complex. The oil

bore through the crank pin was not centered, creating a non standard cross section. The

outer diameter of the crankpin of the ductile cast iron crankshaft was 3.30 cm and the

inner diameter was 1.70 cm. In order to calculate the stress in the cast iron crank-pin,

first the vertical distance to the centroid was found to be 1.57 cm from the bottom of the

cross section shown in Figure 3.23. The area moment of inertia, I, for the cross section

was then calculated using the parallel axis theorem to be 5.29 cm4. The flexure formula,

108

given by Equation 4.1, could then be used to analytically determine the nominal stress at

the critical location. Due to the non-standard cross-section the stress on the top of the

crankpin differed from the stress on the underside of the crankpin. Since the critical

section was determined to be on the underside of the crankpin, the distance from the

centroid to the underside of the crankpin of the cast iron crankshaft, which was 1.57 cm,

was used for c.

The analytical stress results for the forged steel and ductile cast iron crankshafts at

the critical location, which was the crankpin fillet, and without consideration of stress

concentration caused by the fillet, are shown in Table 4.1 (with the locations identified in

Figure 4.2) along with the FEA results from Montazersadgh [2007].

4.2 Finite Element Modeling and Analysis

Finite element modeling was preformed on the forged steel and the ductile cast

iron crankshafts. A linear elastic analysis was used due to the high cycle fatigue situation

which requires nominally elastic loading. The finite element analysis (FEA) was used to

determine the critical location of the crankshafts, determine the stress concentration

factors for the critical fillet locations, and determine the nominal stress applied for the

purpose of life prediction. Using the stress concentration factor approach rather than

directly using the local stress and strain at the critical locations from FEA eliminates the

need for a separate analysis for each load level.

The FEA analysis also validated what was revealed from the experimental stress

results obtained from strain gages. Based on simple analytical calculations of bending

stress (i.e. Mc/I), the stress in the forged steel crankshaft on the top and bottom of the

crankpin should be equal in magnitude due to the symmetry of the crankpin cross-section.

109

The stress results obtained from strain gage readings, however, revealed that there was a

larger stress at the bottom than at the top. This result was confirmed with FEA. The

relatively complex geometry of the crankshaft results in stresses that are not easily

calculated with simple analytical techniques. Finite element analysis was necessary in

this instance in order to account for the complex geometry.

Two types of analyses were performed. First the crankshafts were modeled

according to the dynamic load analysis to determine the critical location of each

crankshaft under in service loading. Second, boundary conditions resembling those of

the test fixture were used in order to determine the stresses for the purpose of life

prediction and to compare with the experimental stress results from strain gages. Details

of the finite element modeling including geometry generation, meshing, boundary

conditions, and loading are presented in Montazersadgh and Fatemi [2007], and

Montazersadgh [2007]. The results relevant to this study are presented in the subsequent

sections.

4.2.1 Critical locations

The critical location is the location of the crankshaft subjected to the highest

stress and therefore the location where fatigue cracks initiate and ultimately lead to

failure. Identification of the critical location was necessary before the component fatigue

testing described in Chapter 3 could be started since the design of the component fatigue

test fixture was based on the location of the critical location. It was expected that the

critical location would be in one of the fillets due to the high stress concentrations at

these locations. The stress contour provided by the FEA based on the dynamic analysis is

110

shown in Figure 4.1. Based on the graphical representation of the stresses, several

locations were selected as potential critical locations of the crankshaft, which are labeled

as locations 1 through 6 in Figure 4.2. The analysis of the critical section is presented for

the forged steel crankshaft, however similar results were obtained for the ductile cast iron

crankshaft.

A plot of the von Mises stresses as a function of crank angle for the locations in

Figure 4.2 are shown in Figure 4.3. The figure shows that a crank angle of 355 degrees

represents the position where the highest stress levels occur. At the crank angle of 355

degrees it can be seen that location 2 identified in Figure 4.2 is the highest stressed

location of the crankshaft. A plot of minimum and maximum stresses, stress range and

mean stress for the critical locations is shown in Figure 4.4. The figure shows that

location number 2, which was determined to have the highest von Mises stress, is also the

location with the highest stress range and mean stress. In fatigue analysis the stress range

and mean stress can be more important than the maximum stress.

Based on the plot of von Mises stress and the stress range and mean stress plot,

location 2 was identified as the critical location of the crankshaft. This location was in

agreement with the literature which has shown that the crankpin fillets are the highest

stressed locations in a crankshaft [Jensen, 1970]. The critical location defined by the

FEA analysis was verified during the component fatigue testing by the fact that all of the

failed crankshafts developed cracks in this crankpin fillet location which then grew to a

large crack which ultimately led to failure.

111

4.2.2 Comparison between FEA, analytical, and experimental results

The boundary conditions in the finite element model were changed from the

dynamic loading condition to resemble the component test assembly. This was done to

investigate the stresses in the component as a result of being loaded in the test apparatus.

As mentioned previously, based on analytical bending stress calculations for the forged

steel crankshaft, the stress at the top and bottom of the crankpin should be symmetric.

However, both experimental and FEA results revealed that the stresses were not

symmetric. The comparison between FEA and experimental values was also done to

validate the finite element model. The experimental results were only obtained and

compared to the FEA results for the forged steel crankshaft. It was not necessary to

repeat the procedure for the cast iron crankshaft since it was modeled and tested the same

way as the forged steel crankshaft. The details of the finite element analysis are given in

Montazersadgh and Fatemi [2007].

Strain gages were mounted at the four locations labeled a, b, c, and d in Figure

4.2. The crankshaft was installed in the test fixture such that locations a and b were on

the top and bottom respectively. The front main bearing was clamped in the support arm

and the rear main bearing (the right side of Figure 4.2) was clamped in the moment arm.

It should be noted that this was not the position the crankshafts were tested in as this

comparison was done prior to properly identifying the critical fillet. In the tests, as

described in Chapter 3, the crankshaft was switched such that the rear main bearing was

in the support arm and moment arm was clamped to the front main bearing. Despite this

change, the results were applicable since they were used for comparison purposes. The

load was applied vertically through the moment arm attached to the right of Figure 4.2.

112

The crankshaft was also rotated 90 degrees so that locations c and d (Figure 4.2) were on

the top and bottom, respectively. This was done to see the effects of the different offsets

of the crank-pin.

The results from FEA, analytical calculations, and experimental results from

strain gages are shown in Table 4.2. The table shows good agreement between the

experimental results obtained from strain gages and the results from FEA. All of the

differences between FEA and experimental results were 6.5% or less. The results from

the analytical analysis are also close to the FEA and experimental results considering the

complex geometry.

The analytical results, discussed in Section 4.2, suggest that the stress at location

a and location b should be equal and opposite. FEA is necessary in order to obtain

accurate stresses in the component which, in a case where the geometry is complex, could

otherwise not be obtained. Both FEA and experimental results show that the magnitudes

of the stresses at locations a and b (identified in Figure 4.2) are not equal due to the

complexity of the geometry. However, when the crankshaft was rotated 90 degrees, the

FEA results show that stresses at location c and location d are equal in magnitude, which

is confirmed by the experimental results from the strain gages. Some difference is seen in

the strain gage readings which can be attributed to errors in placing the gages.

4.2.3 FEA results used for life predictions

The finite element analysis with the boundary conditions resembling the

component test fixture was used to determine the nominal stresses in the crankshafts.

This nominal stress was used for life prediction purposes. For this analysis a point load

113

of 4.45kN was applied at B, as shown in Figure 4.2 for the forged steel crankshaft, which

was the same for the cast iron crankshaft. The FEA results of the test condition are

shown in Table 4.3 for forged steel and Table 4.4 for the cast iron crankshaft where the

locations are identified in Figure 4.2. The finite element analysis based on the

component test fixture showed that the highest stress was at the previously identified

critical section, indicating that the test set-up would produce failures in the critical

section. The stress at location b as shown in Figures 4.2 was extrapolated to determine

the nominal stress at the critical location, point 2 in Figure 4.2. For the forged steel

crankshaft the moment arm to location 4 was 12.6 cm resulting in a nominal stress of

121.6 MPa when 4.45 kN force was applied. To obtain the nominal stress at the critical

location a ratio was used to accommodate a longer moment arm and different loads

corresponding to the actual tests. The ratio related the stress from FEA, which was

obtained from a given bending moment, to the stress resulting from a different bending

moment. This was possible because of the uniform cross-section and the stress being

linearly related to the bending moment. A similar procedure was used for the ductile cast

iron crankshaft.

The nominal stress values at the critical fillet, location b in Figure 4.2, obtained

from the FEA [Montazersadgh, 2007] for the forged steel and ductile cast iron

crankshafts are shown in Table 4.1 for the different applied moments. The results were

used in the life predictions discussed in Sections 4.3 and 4.4.

Determining the stress concentration factor (Kt) was also necessary for life

predictions. The stress concentration factor allowed the stress at the fillets to be

calculated knowing the nominal stress which was obtained from the linear elastic finite

114

element analysis. As shown in Table 4.3 for forged steel, the stress obtained from FEA

with a 4.45 kN applied load was 121.6 MPa for location b and 539.7 MPa for location 2

(locations identified in Figure 4.2). Location b was considered to be far enough away

from fillets such that it was not affected by stress concentrations. Therefore, as was done

above, this stress was used to determine the nominal stress at location 2 by equating the

ratio of the stresses to the ratio of the moment arms. The distance from the applied load

to location b was 12.6 cm and to location 2 was 14.3 cm. The nominal stress at the

critical location was determined to be 137.1 MPa. The stress concentration factor, Kt,

was calculated by dividing the stress at the critical location from FEA by the nominal

stress at the critical location:

nomt S

SK = (4.2)

The stress concentration factor for forged steel was determined to be 3.94. A similar

procedure was followed for the cast iron crankshaft where the nominal stress at the

critical location was determined to be 160.7 MPa and the stress concentration factor, Kt,

was 3.09.

4.3 Stress-Life Approach and Life Predictions

4.3.1 Procedures and predictions

In service the crankshaft is subjected to very high cycle fatigue, requiring the stresses

to be elastic. In situations where stresses are predominately elastic and high cycle fatigue

is present, the stress-life (S-N) approach is commonly used. The S-N approach uses the

nominal stress rather than the localized stress at the root of the notch. To account for the

115

stress concentration at the notch, the fatigue limit is reduced by the fatigue notch factor.

The effect of mean stress can be accounted for by an equation such as the modified

Goodman equation. The S-N line can also be modified to account for other effects such

as surface finish effect.

As a starting point, the stress-life curve for a smooth, unnotched member,

subjected to completely reversed loading, was constructed from the fatigue strength

coefficient, σf’, and the fatigue strength exponent, b. Both σf’ and b were obtained from

the specimen tests as presented in Chapter 2 and the values for both materials are listed in

Table 2.4. The S-N line for the smooth, unnotched member is represented by the

equation:

( ) bffS 6' 102×= σ (4.3)

The effect of the notch was taken into account by the fatigue notch factor, Kf .

The fatigue notch factor depends on the geometry of the notch and also notch sensitivity

of the material. The notch sensitivity of a material is defined by:

11

−

−=

t

f

KK

q (4.4)

where Kf is the fatigue notch factor and Kt is the stress concentration factor. A value for q

= 0 indicates no sensitivity to notches and q = 1 is defined as full notch sensitivity. When

a material has large sensitivity to notches Kf is approximately equal to Kt. There are

several equations for estimating the notch sensitivity of a material, including Peterson’s

equation given by [Stephens et al., 2000]:

raq

+=

11 (4.5)

116

where q is the notch sensitivity, r is the radius of the notch and a is the material

characteristic length. The following equation was used to calculate the material

characteristic length, a (in millimeters) for the forged steel material [Stephens et al.,

2000]:

( ) 8.120700254.0uSa = (4.6)

where Su is the ultimate strength of the material in MPa. The notch sensitivity for the

forged steel was calculated to be 0.95, indicating very high notch sensitivity. The fatigue

notch factor for the forged steel crankshaft with a notch radius of r = 2.38 mm was then

calculated to be 3.78.

For the ductile cast iron crankshafts, Peterson’s equation (Equation 4.5) was also

used along with the equation for characteristic length, a, intended for steels (Equation

4.6). Using this approach the notch sensitivity was calculated to be 0.92, which for the

ductile cast iron crankshaft with a fillet radius of r = 2.45 mm resulted in Kf = 2.93. The

life predictions using this approach were conservative. The difference in stress

concentrations between the forged steel and ductile cast iron crankshafts can be attributed

to the geometric differences in the two crankshafts. The ductile cast iron crankshaft has a

difference in stiffness in the web section in close proximity to the critical fillet, as

compared to the forged steel crankshaft, which decreases the stress in the fillet, thus

decreasing the stress concentration.

According to Shigley and Mitschke [2002], cast iron has very low notch

sensitivity, q, ranging in value from 0, or no notch sensitivity, to 0.2. Juvinall and

Marshek [1991] also state that cast irons have little to no notch sensitivity. Cast iron has

inclusions and porosity which can act as notches. Therefore, these notch effects are

117

already taken into account in the cyclic material properties obtain from the strain-

controlled specimen fatigue testing. Shigley and Mitschke recommend that to be

conservative, a value of q = 0.2 be used for all cast irons [2002]. Using q = 0.2 resulted

in Kf = 1.42. The results of the life predictions were excessively nonconservative when

this approach was used to calculate Kf. The true value for notch sensitivity, therefore, lies

somewhere between the values obtained using these two approaches. The life predictions

are presented using both values of Kf, however, the approach using the higher fatigue

notch factor, Kf = 2.93, is considered the primary approach due to the better agreement

with the component fatigue tests data.

The fatigue life at 2 x 106 reversals was reduced by Kf. Therefore, for the notched

member, one point on the S-N line was Sf /Kf at 2 x 106 reversals. The forged steel

crankshaft had a notched fatigue strength at 2 x 106 reversals of 94.4 MPa. The ductile

cast iron is also assumed to have a fatigue limit at 2 x 106 reversals [Juvinall and

Marshek, 1991]. The notched fatigue strength at 2 x 106 reversals for the cast iron

crankshaft based on Kf = 2.93 was 89.5 MPa, and based on Kf = 1.42 the notched fatigue

strength was 185.0 MPa.

The point at 2 x 106 reversals was connected to the stress amplitude, σf’, at one

reversal for each material. This approach assumes that there is no effect of the notch at

one cycle due to the presence of gross plastic deformation, or yielding.

The notched S-N line for the forged steel crankshaft is represented by:

( ) 1704.021124 −= ff NS (4.7)

and the notched S-N line for the ductile cast iron crankshaft with Kf = 2.93 is represented

by:

118

( ) 1611.02927 −= ff NS (4.8)

The S-N lines obtained above assume R = -1. To account for mean stress which is

present in situations where the loading is not completely reversed, the modified Goodman

equation is often used. The modified Goodman equation is given by [Stephens et al.,

2000]:

1=+u

ma

SSS

fKNfS (4.9)

where Sa is the alternating stress, Sm is the mean stress, Su is the ultimate strength of the

material, and SNf is the fully reversed fatigue strength at 2 x 106 reversals.

Dynamic load analysis determined the stress ratio, R, of the crankshafts to be

approximately -0.2. The details of the dynamic load analysis are shown in Montazersadgh

and Fatemi [2007] and Montazersadgh [2007]. The R-ratio was used to calculate the

mean stress. The R-ratio is defined by:

max

min

SS

R = (4.10)

For R = -0.2 the relationship between Sa and Sm was calculated to be:

am SS 667.0= (4.11)

The equation for the notched R = -0.2 condition for the forged steel crankshaft is

represented by:

( ) 906.021124

170.0 +=

fa N

MPaS (4.12)

which results in a fatigue strength of 87.8 MPa at 2 x 106 reversals. The equation for the

notched R = -0.2 condition and high notch sensitivity (Kf = 2.93) for the ductile cast iron

crankshaft is then represented by:

119

( ) 939.02927

161.0 +=

fa N

MPaS (4.13)

which results in a fatigue strength of 82.1 MPa at 2 x 106 reversals. For the low notch

sensitivity (Kf = 1.42) assumption of ductile cast iron, the S-N curve is represented by:

( ) 939.02927

111.0 +=

fa N

MPaS (4.14)

which results in a fatigue strength of 155.8 MPa at 2 x 106 reversals.

The unnotched R = -1, notched R = -1, and notched R = -0.2 S-N curves are

shown in Figure 4.5 for the forged steel crankshaft and in Figure 4.6 for the ductile cast

iron crankshaft. Results of the S-N life predictions for the forged steel and for the ductile

cast iron crankshafts are shown in Tables 4.5 and 4.6, respectively.

4.3.2 Comparisons with experimental results

The results of the S-N life predictions along with the crankshaft fatigue test data

are summarized in Table 4.7 for the forged steel and ductile cast iron crankshafts. The

forged steel predicted S-N line superimposed with the experimental data points is shown

in Figure 4.7 for the crack initiation criterion and in Figure 4.8 for the 5% change in

displacement amplitude criterion. The figures show a very good agreement between the

predictions and the experimental data for both failure criteria.

The ductile cast iron crankshaft predicted S-N lines, using both values of Kf,

superimposed with the experimental data points are shown in Figure 4.9 for the crack

initiation criterion and in Figure 4.10 for the 5% change in displacement amplitude

criterion. When Kf = 1.42 is used, the S-N life predictions for both the crack initiation

and 5% displacement amplitude are nonconservative, while the opposite is true if the

120

value for Kf = 2.93 is used. Based on the S-N lines, the predictions using Kf = 2.93 more

closely match the experiment data than when Kf = 1.42 was used. For the S-N approach

for the ductile cast iron crankshafts, the most accurate predictions are obtained using Kf =

2.93 and the crack initiation failure criterion as evidenced by the close proximity of the

experimental data points to the predicted S-N curve in Figure 4.9. Since the predictions

using the low notch sensitivity assumed for cast iron (Kf = 1.42) were not in agreement

with the experimental data, further comparisons using this assumption are not presented.

The comparison between experimental results and predicted results are shown in

Figure 4.11 for the crack initiation failure criterion for both the forged steel and ductile

cast iron crankshafts (Kf = 2.93). In Figure 4.11, the center line with a slope of one (45

degrees) and passing through the origin represents a perfect correlation between the

prediction and experimental data. Data points that are above the line represent an over

prediction (non-conservative), while points below the line represent an under prediction

(conservative). The other lines represent factors of two and three differences. The

experimental results versus predicted results are plotted for the 5% change in

displacement amplitude criterion in Figure 4.12 for both crankshafts.

Figures 4.11 and 4.12 show that for the forged steel crankshafts, the experimental

data fall within a factor of two of the prediction for both the crack initiation and 5%

change in displacement amplitude criteria. Comparison of Figures 4.11 and 4.12 reveals

that the crack initiation data is in better agreement with the prediction than the 5% change

in displacement amplitude criterion. This is expected, as the failure for specimen fatigue

tests was based on crack initiation, and data from these tests were used for crankshaft life

121

predictions. The prediction is more conservative when the 5% change in displacement

amplitude criterion is used.

The predictions for the cast iron crankshaft were less in agreement with the

predictions for the forged steel crankshaft although the predictions were still reasonably

accurate when the higher notch sensitivity was used. The experimental data for the cast

iron crankshaft are along the factor of two and factor of three scatter bands when the

crack initiation failure criterion is used as shown in Figure 4.11. When the 5% change in

displacement amplitude criterion is used, the experimental data are along or slightly

outside of the factor of three scatter band as shown in Figure 4.12 with predictions being

conservative.

The S-N prediction took into account the stress concentration caused by the

crankpin fillet and the mean stress effect. However, the effect of surface finish was

neglected for both the forged steel and cast iron crankshafts. The surface of the both

crankshafts appeared to be ground and, therefore, a very smooth surface finish with few

machining marks, which was approximated as a smooth surface. Had the critical location

been in an area that was in the as forged or as cast condition, a surface finish correction

factor would have been needed.

The fillets of crankshafts in automotive applications are often rolled to induce

compressive residual stresses. The residual stress will, in a long life situation, provide

better fatigue performance. However, residual stresses were not considered since the

crankshafts were not rolled in this case.

122

4.4 Stain-Life Approach and Life Predictions

4.4.1 Procedures and predictions

The strain-life approach to life estimation is commonly used in low cycle fatigue

applications due to the presence of inelastic strain in the component. In addition, strains

can be measured in complex geometries and at stress concentrations, resulting in an easy

comparison with data obtained from strain-controlled specimen fatigue tests. Although

crankshafts are a high cycle fatigue component, the strain-life approach can still be

valuable due the presence of notches in the crankshaft. The strain-life approach is

commonly used for notched members, because local plastic deformation often occurs at

the root of the notch, even when an elastic loading condition is present. This approach

uses the stresses and strains at the root of the notch, as opposed to the S-N approach

which uses nominal stresses. In the stress-life approach the nominal stresses are known

and, therefore, the life to crack initiation can be directly calculated. However, in the

strain-life approach, first the notch stresses and strains must be determined.

There are several methods which can be used to calculate the local stress and

strain at the root of the notch given the nominal elastic stress. Analytical methods to

calculate the local stress and strain include the linear rule, Neuber’s rule, and Glinka’s

rule. Neuber’s rule, the most commonly used model, is presented here. In the case of a

plane strain situation, Glinka’s rule is more applicable [Stephens et al., 2000]. For

comparison the notch stresses and strains were also calculated using Glinka’s rule and the

results were very similar to those obtained using Neuber’s rule.

Neuber’s rule assumes that the geometric mean of the stress concentration and

strain concentration factors remain constant under plastic deformation and are also equal

123

to the stress concentration factor. Graphically, the notch stress and strain are determined

from the intersection of the stress-strain curve and the Neuber hyperbola. The stress-

strain curve is represented by the Ramberg-Osgood equation given by:

npe KE

1

⎟⎠⎞

⎜⎝⎛+=+=

σσεεε (4.15)

while Neuber’s hyperbola is represented by:

eSKt2=σε (4.16)

where S and e are the nominal stress and strains, and σ and ε are the stress and strain at

the root of the notch. Therefore, the intersection point can be found by solving equations

4.15 and 4.16 simultaneously.

The nominal stress is typically elastic, otherwise there will be gross plastic

deformation and the part fails by yielding rather than fatigue crack initiation. In the case

where the nominal stress is elastic, the engineering strain, e, is simply the nominal stress

divided by modulus of elasticity (S/E). Therefore, Neuber’s rule for nominal elastic

behavior becomes:

( )ESKt

2

=σε (4.17)

In the case of cyclic loading, which is the case when using Neuber’s rule for

fatigue life predictions, the stress-strain curve is replaced with the stable hysteresis loop

represented by the equation [Stephens et al., 2000]:

'1

'22

n

KE ⎟⎟⎠

⎞⎜⎜⎝

⎛ ∆+

∆=∆

σσε (4.18)

Equation 4.18 assumes that the material exhibits a Massing behavior, with a factor of 2,

meaning that the hysteresis loop can be obtained by doubling the cyclic stress-strain

124

curve. For cyclic loading, the stresses and strains are replaced with the stress and strain

ranges and Neuber’s rule becomes:

( )E

SK f2∆

=∆∆ σε (4.19)

It should be noted that in this equation Kt was replaced with Kf which has been shown by

Topper et al. [Stephens et al., 2000] to have better agreement with experimental data. For

the first cycle ∆ε, ∆σ, and ∆S in equations 4.18 and 4.19 were replaced with εmax, σmax,

and Smax, respectively. This approach assumes that the maximum stress is reached in the

first cycle, whereas in testing the load gradually increased to the maximum values over

approximately 100 cycles. Therefore, there could be some differences in lives between

using this approach versus performing the analysis based on the test gradually reaching

the maximum load.

After the notch stresses and strains are determined, the life to crack initiation can

be predicted. The prediction estimates the life to the onset of a crack on the order of 1

mm. Several equations exist for predicting fatigue life in the presence of mean stresses,

based on the strain-life approach, including Morrow’s mean stress parameter, and the

Smith–Watson–Topper (SWT) parameter. Both equations account for the mean stress

effects. The SWT parameter has been shown to be applicable to a broad range of

materials [Stephens et al., 2000]. Due to its broad applicability, the SWT parameter was

used in this study.

The SWT parameter assumes that the product of maximum notch stress and notch

strain amplitude (σmax εa) remains constant regardless of the individual contribution of

notch strain amplitude, εa, and notch mean stress σm. The SWT equation is represented

by:

125

( ) ( ) ( ) cbfff

bffa NENE ++= 2''2' 22

max εσσεσ (4.20)

Using the material properties for each material and the product of σmax and εa for each test

level which were determined using Neuber’s rule, the fatigue life was predicted using

Equation 4.20.

4.4.2 Comparisons with experimental results

The SWT parameter versus reversals to failure for the forged steel crankshaft is

shown in Figure 4.13 for the crack initiation criterion and in Figure 4.14 for the 5%

change in displacement amplitude criterion. The same plots are shown for the cast iron

crankshaft in Figures 4.15 and 4.16. For the cast iron crankshafts only predictions for the

higher notch sensitivity (Kf = 2.93) are presented due to the S-N approach showing that

the low notch sensitivity predictions were not in agreement with experimental data. The

results from the strain-life predictions using the SWT parameter along with the notch

stresses and strains obtained from Neuber’s rule are summarized in Table 4.5 for the

forged steel and in Table 4.6 for the cast iron crankshafts. The predictions along with the

crankshaft experimental data are also presented in Table 4.7.

For the forged steel crankshaft, the strain life approach resulted in predictions that

reasonably agreed with the experimental data when the crack initiation failure criterion

was used, as shown in Figure 4.13. When the 5% change in displacement amplitude

failure criterion was used the predictions were also reasonable, as shown in Figure 4.14.

The predictions, however, more closely agreed with the data using the crack initiation

criterion.

126

For the cast iron crankshafts, where the high notch sensitivity was assumed, the

strain-life predictions under-estimated the fatigue lives. This was true for both the crack

initiation failure criterion as shown in Figure 4.15, and the 5% change in displacement

amplitude criterion as shown in Figure 4.16. Although the predictions were less accurate

than they were for the forged steel crankshaft, the predictions were always conservative.

The predicted cycles to failure using the strain-life approach versus experimental

cycles to failure using the crack initiation criterion are shown in Figure 4.17 for the

forged steel and ductile cast iron crankshafts, and in Figure 4.18 for the 5% change in

displacement amplitude criterion. Figure 4.17 shows that for the forged steel crankshaft,

the predictions were reasonably accurate as all of the data points are inside the factor of 3

scatter band based on the crack initiation criterion. Figure 4.18 shows that for the forged

steel crankshaft the predictions were more conservative when compared to the 5% change

in displacement experimental data. For the cast iron crankshaft, all of the data points

were outside of the factor of 3 scatter band, indicating less accurate, although

conservative predictions, when compared to experimental data based on both failure

criteria. For both the forged steel and ductile cast iron crankshafts, the strain-life

predictions were more accurate when compared to the crack initiation failure criterion.

4.5 Discussion of Life Prediction Results

The results of the predictions using the stress-life and strain-life are shown in

Table 4.7 along with the component test data. The strain-life approach in this case

resulted in shorter fatigue life predictions than the S-N approach, which resulted in the

127

strain-life approach always under-predicting the fatigue lives of both the forged steel and

ductile cast iron crankshafts. Therefore, the strain-life approach always provided

conservative fatigue life predictions. Both methods of predictions resulted in more

accurate comparisons for the forged steel crankshaft than the ductile cast iron crankshaft.

For the forged steel crankshaft, life predictions using the stress-life and strain-life

approaches were reasonable for both the crack initiation and 5% change in displacement

amplitude criteria. However, both the S-N and strain-life approaches were more accurate

when the crack-initiation failure criterion was used, compared to the 5% change in

displacement amplitude criterion. This result is reasonable since the fatigue life for

prediction purposes is considered to be the life to the onset of cracks on the order of a

millimeter. By the time the change in displacement amplitude was 5%, the crack was

already much longer than 1 mm. Therefore, it was expected that the crack-initiation data

would better fit the life predictions. The strain-life approach resulted in predictions that

were conservative, while the S-N approach did not always result in a conservative

predictions. The S-N approach predicted longer lives than the strain-life approach, and in

this case the S-N approach sometimes predicted longer lives than what was observed in

the crankshaft fatigue tests. The strain-life approach always predicted lives that were less

than the actual fatigue lives of the crankshafts when compared to experimental data,

making the strain-life approach more conservative.

128

Table 4.1: Analytical nominal stress results at the critical location and comparison with FEA results for the forged steel and cast iron crankshafts.

Forged Steel Stress (MPa) Cast Iron Stress (MPa)

Analytical FEA Analytical FEA Moment

Amplitude (N-m)

Location a and b

Location a

Location b

Location a

Location b

Location a

Location b

630 134.7 118.9 140.3 206.3 187.5 218.9 185.6

517 110.4 97.6 115.0 169.1 153.7 179.6 152.1

431 92.0 81.4 95.9 141.0 128.1 149.8 126.8

350 74.8 66.1 78.0 114.6 104.2 121.6 103.1

129

Table 4.2: Comparison between FEA, experimental, and analytical stress results for the forged steel crankshaft in the as-tested condition at the locations shown in Figure 4.2.

At Location a

Load (N)

FEA (MPa)

EXP (MPa)

% Difference between FEA

and EXP ANALYTICAL

(MPa) -890 -61.6 -59.3 3.80% -72.4 890 61.5 65.5 6.50% 72.4

At Location b

Load (N)

FEA (MPa)

EXP (MPa)

% Difference between FEA

and EXP ANALYTICAL

(MPa) -890 86.9 81.4 6.30% 72.4 890 -86.7 -90.3 4.20% -72.4

At Location c

Load (N)

FEA (MPa)

EXP (MPa)

% Difference between FEA

and EXP ANALYTICAL

(MPa) -890 -76.4 -71.7 6.10% -72.4 890 76.3 75.8 0.50% 72.4

At Location d

Load (N)

FEA (MPa)

EXP (MPa)

% Difference between FEA

and EXP ANALYTICAL

(MPa) -890 75.5 71.7 5.00% 72.4 890 -75.6 -76.5 1.30% -72.4

130

Table 4.3: FEA results for the test setup boundary conditions for the forged steel crankshaft for the locations identified in Figure 4.2.

Location Load (kN)

Stress (MPa)

Moment Arm (cm)

1 405.08 -- 2 539.74 14.26 3 374.74 11.04 4 52.52 14.26 5 76.53 11.04 6 161.48 -- 8 155.90 -- 9 392.85 -- a 106.22 12.65 b

4.45

121.63 12.65

Table 4.4: FEA results for the test setup boundary conditions for the cast iron crankshaft

for the locations identified in Figure 4.2.

Location Load (kN)

Stress (MPa)

Moment Arm (cm)

2 496.86 12.29 a 159.77 10.34 b

4.45 135.31 10.34

131

Table 4.5: Life prediction results including the S-N and ε-N approaches for the forged steel crankshaft.

S-N ε-N: Neuber's Rule Moment Amplitude

(N-m)

Sa (MPa) Predicted Nf

∆σ (MPa) ∆ε εaσmax Predicted Nf

630 140.3 49,695 916.2 0.0056 1.580 23,163

517 115.0 182,710 811.1 0.0042 1.116 74,074

431 95.9 585,600 705.0 0.0017 0.828 250,786

350 78.0 >106 584.7 0.0037 0.592 >106

Table 4.6: Life prediction results including the S-N and ε-N approaches for the ductile cast iron crankshaft.

S-N ε-N: Neuber's Rule

Kf = 2.93 Kf = 1.42 Kf = 2.93 Moment

Amplitude (N-m)

Sa (MPa)

Predicted Nf ∆σ

(MPa) ∆ε εaσmax Predicted Nf

630 185.6 2,978 149,321 947.2 0.0070 2.022 691

517 152.1 13,219 >106 840.0 0.0053 1.435 1,946

431 126.8 49,125 >106 728.2 0.0043 1.072 5,754

350 103.1 210,216 >106 601.3 0.0034 0.774 24,961

132

Table 4.7: Experimental data and life prediction results for the forged steel and ductile cast iron crankshafts.

Applied Moment Amp. (N-m) Crack Initiation Cycles at 5% Change

in Disp. Amp. S-N Prediction ε-N Prediction

Forged Steel Crankshaft 29,248 45,568 49,695 23,163 45,302 69,670 49,695 23,163 630 58,236 90,853 49,695 23,163

145,000 234,289 182,710 74,074 98,741 213,885 182,710 74,074 517 204,174 396,011 182,710 74,074

>2.09 x 106 >2.09 x 106 585,600 250,786 431 >3.980 x 106 >3.980 x 106 585,600 250,786

350 >3.24 x 106 >3.24 x 106 >106 >106 Cast Iron Crankshaft

Kf = 1.42 Kf = 2.93 Kf = 2.93 7,132 17,353 149,321 2,978 691 9,256 17,380 149,321 2,978 691 630 8,021 20,957 149,321 2,978 691

25,512 47,513 >106 13,219 1,946 24,096 52,790 >106 13,219 1,946 517 37,380 54,966 >106 13,219 1,946 75,200 132,877 >106 49,125 5,754 78,367 121,866 >106 49,125 5,754 431 82,200 143,259 >106 49,125 5,754

920,783 1,005,665 >106 210,216 24,900 350 301,774 370,216 >106 210,216 24,900

133

Figure 4.2: Forged steel crankshaft showing the analyzed locations for the dynamic load analysis and dynamic based FEA.

98

B

Figure 4.1: Forged steel crankshaft showing FEA stress contour with the crankpin fillet magnified [Montazersadgh, 2007]

134

-50

0

50

100

150

200

0 180 360 540 720

Crankshaft Angle (Deg)

Stre

ss M

agni

tude

(MPa

)

1 2 3 4 5 6

Figure 4.3: Stress magnitude versus crankshaft angle for the locations shown in Figure 4.2 [Montazersadgh and Fatemi, 2007].

-50

0

50

100

150

200

250

1 2 3 4 5 6

Location Number

Stre

ss M

agni

tude

(MP

a)

Maximum Minimum Range Mean

Figure 4.4: Maximum stress, minimum stress, stress range, and mean stress results from FEA for the locations shown in Figure 4.2 [Montazersadgh and Fatemi, 2007].

135

10

100

1000

10000

1.E+00 1.E+01 1.E+02 1.E+03 1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

Stre

ss A

mpl

itude

(MPa

)

Unnotched, R = -1Notched, R = -1Notched, R = -0.2

Figure 4.5: Forged steel crankshaft S-N lines for the unnotched, notched, and notched R = -0.2 condition.

10

100

1000

10000

1.E+00 1.E+01 1.E+02 1.E+03 1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

Stre

ss A

mpl

itude

(MPa

)

Unnotched, R = -1

Notched, R = -1

Notched, R = -0.2, Kf=2.93

Notched, R = -0.2, Kf=1.42

Kf = 2.93

Kf = 1.42

Figure 4.6: Ductile cast iron crankshaft S-N lines for the unnotched, notched, and notched R = -0.2 condition.

136

10

100

1000

1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

Stre

ss A

mpl

itdue

(MPa

)

Notched, R = -0.2 Prediction

Crack Initiation Experimental Data

Figure 4.7: Forged steel crankshaft S-N line for the notched R = -0.2 condition superimposed with the crack initiation experimental data.

10

100

1000

1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

Stre

ss A

mpl

itdue

(MPa

)

Notched, R = -0.2 Prediction

5% Change in Disp. Amp. ExperimentalData

Figure 4.8: Forged steel crankshaft S-N line for the notched R = -0.2 condition superimposed with the 5% change in displacement amplitude experimental data.

137

10

100

1000

1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

Stre

ss A

mpl

itude

(MPa

)

Cast Iron, Notched R = -0.2, Kf=2.93

Cast Iron, Notched R = -0.2, Kf=1.42

Crack Initiation Experimental Data

Kf = 2.93

Kf = 1.42

Figure 4.9: Ductile cast iron crankshaft S-N lines for the notched R = -0.2 condition

superimposed with the crack initiation experimental data.

10

100

1000

1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

Stre

ss A

mpl

itude

(MPa

)

Cast Iron, Notched R = -0.2, Kf=2.93

Cast Iron, Notched R = -0.2, Kf=1.42

5% Change in Disp. Amp. Experimental Data

Kf = 2.93

Kf = 1.42

Figure 4.10: Ductile cast iron crankshaft S-N lines for the notched R = -0.2 condition

superimposed with the 5% change in displacement amplitude experimental data.

138

1.E+03

1.E+04

1.E+05

1.E+06

1.E+07

1.E+03 1.E+04 1.E+05 1.E+06 1.E+07

Experimental Cycles to Failure

Pred

icte

d C

ycle

s to

Failu

re

Forged Steel

Cast Iron Kf = 2.93

Figure 4.11: Predicted versus experimental cycles to failure using the S-N approach for

the forged steel and ductile cast iron crankshafts using the crack initiation failure criterion.

1.E+03

1.E+04

1.E+05

1.E+06

1.E+07

1.E+03 1.E+04 1.E+05 1.E+06 1.E+07

Experimental Cycles to Failure

Pred

icte

d C

ycle

s to

Failu

re

Forged Steel

Cast Iron

(3)

Kf = 2.93

Figure 4.12: Predicted versus experimental cycles to failure using the S-N approach for

the forged steel and ductile cast iron crankshafts using the 5% change in displacement amplitude failure criterion.

139

0.1

1

10

1.E+04 1.E+05 1.E+06 1.E+07Reversals to Failure (2Nf)

SWT

Par

amet

er (σ

maxε a

)

Forged Steel Predicted

Forged Steel Data - Neuber

Figure 4.13: SWT parameter versus reversals to failure based on crack initiation with

strain-life prediction data superimposed for the forged steel crankshafts.

0.1

1

10

1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

SWT

Par

amet

er (σ

maxε a

)

Forged Steel Predicted

Forged Steel Data - Neuber

Figure 4.14: SWT parameter versus reversals to failure based on 5% change in

displacement amplitude with strain-life prediction data superimposed for the forged steel crankshafts.

140

0.1

1

10

1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

SWT

Par

amet

er (σ

maxε a

)

Cast Iron Predicted

Cast Iron Data Kf = 2.93

Figure 4.15: SWT parameter versus reversals to failure based on crack initiation with

strain-life prediction data superimposed for the ductile cast iron crankshafts.

0.1

1

10

1.E+04 1.E+05 1.E+06 1.E+07

Reversals to Failure (2Nf)

SWT

Par

amet

er (σ

maxε a

)

Cast Iron Predicted

Cast Iron Data Kf = 2.93

Figure 4.16: SWT parameter versus reversals to failure based on 5% change in

displacement amplitude with strain-life prediction data superimposed for the ductile cast iron crankshafts.

141

1.E+02

1.E+03

1.E+04

1.E+05

1.E+06

1.E+07

1.E+02 1.E+03 1.E+04 1.E+05 1.E+06 1.E+07Experimental Cycles to Failure

Pred

icte

d C

ycle

s to

Failu

re

Forged Steel

Cast Iron Kf = 2.93

Figure 4.17: Predicted versus experimental cycles to failure using the strain-life

approach for the forged steel and ductile cast iron crankshafts based on the crack initiation failure criterion.

1.E+02

1.E+03

1.E+04

1.E+05

1.E+06

1.E+07

1.E+02 1.E+03 1.E+04 1.E+05 1.E+06 1.E+07

Experimental Cycles to Failure

Pred

icte

d C

ycle

s to

Failu

re

Forged Steel

Cast Iron

(3)

Kf = 2.93

Figure 4.18: Predicted versus experimental cycles to failure using the strain-life

approach for the forged steel and ductile cast iron crankshafts based on the 5% change in displacement amplitude failure criterion.

142

CHAPTER 5

SUMMARY AND CONCLUSIONS

The fatigue behaviors of forged steel and cast iron crankshafts from a one

cylinder engine were obtained and compared. In order to compare the two crankshafts,

first specimen testing was carried out on specimens machined from the as-forged and as-

cast crankshafts. Specimen testing included tensile tests to obtain the monotonic material

properties, strain-controlled uniaxial fatigue tests to obtain the cyclic properties of the

two materials, and Charpy V-notch impact tests to determine the impact toughness of the

materials. Load-controlled bending fatigue tests with R = -0.2 were then performed on

the crankshafts. Results from finite element analysis [Montazersadgh and Fatemi, 2007]

were used to obtain the stresses in the crankshafts. Fatigue life predictions using the S-N

and ε-N approaches were then carried out using the stress results from FEA. Based on

the experimental results and the analyses performed the following conclusions were

drawn:

1. Based on the monotonic tensile test results, the forged steel has significantly higher

strength than the ductile cast iron. The yield strength of the forged steel is 52%

higher than that of the cast iron, while the ultimate strength is 26% higher for the

forged steel than the ductile cast iron.

143

2. The forged steel material also has more ductility than the ductile cast iron as shown

by the percent reduction in area, which was 58% for the forged steel and 6% for the

ductile cast iron.

3. The forged steel Charpy V-notch impact results show that the forged steel in both the

L-T and T-L directions have higher impact toughness than the ductile cast iron at all

temperature levels investigated. This is important for this application due to the

possibility of impact loading condition in the engine if subjected to a sudden stop.

4. The S-N curves for the two materials show that the forged steel has better fatigue

resistance than the ductile cast iron. The fatigue strength at 106 cycles was 359 MPa

for the forged steel and 263 MPa for the ductile cast iron, which results in a factor of

30 longer life for the forged steel in the long life region. The forged steel fatigue

strength at 106 cycles is 36% higher than the ductile cast iron.

5. The forged steel also shows longer life when subjected to plastic deformation, based

on the true plastic strain amplitude versus reversals to failure plot. For a given

plastic strain amplitude, the forged steel has a factor of 40 longer life than the ductile

cast iron.

6. The Neuber curves for the two materials also show better fatigue performance for the

forged steel material, compared to the ductile cast iron. The Neuber curves show

that in the long life region the forged steel has a factor of 50 longer life than the

ductile cast iron.

7. The crack growth life for both crankshafts was a significant portion of the fatigue

life during the crankshaft testing. The crack growth rate of the forged steel

crankshaft was slower than the ductile cast iron crankshaft.

144

8. The failure criterion based on crack initiation is more reasonable in crankshaft

applications since an engine would not tolerate the increased deflection caused by

the presence of a crack. The 5% change in displacement criterion resulted in a crack

that was 10 mm or longer.

9. Based on the crack initiation failure criterion the forged steel crankshaft had a factor

of 6 longer life than the ductile cast iron crankshaft at long lives. The 5% change in

displacement amplitude also showed better fatigue performance for the forged steel

crankshaft, resulting in an order of magnitude longer life than the ductile cast iron

crankshaft at long lives.

10. At 106 cycles the fatigue strength of forged steel crankshaft was 36% higher than the

fatigue strength of the ductile cast iron crankshaft. Specimen fatigue test results also

show that the fatigue strength of the forged steel material was 36% higher than the

fatigue strength of the ductile cast iron material at 106 cycles.

11. During crankshaft fatigue tests, circumferential cracks developed in the rear crankpin

fillet of both forged steel and ductile cast iron crankshafts which was identified as

the critical location from FEA. These cracks grew and were the ultimate cause of

failure for the crankshafts, despite secondary cracks which developed in the opposite

crankpin fillet in some crankshafts.

12. Finite element analysis was necessary to obtain the stresses in the crankshafts due to

the relatively complex geometry. The geometry led to a lack of symmetry at the top

and bottom of the crankpin in the forged steel crankshaft in spite of cross-section

symmetry, which could not be accounted for in the analytical stress calculations.

145

The lack of symmetry at the top and bottom of the crankpin in the forged steel

crankshaft was confirmed with experimental strain gage results.

13. The life predictions were more accurate for the forged steel crankshafts than the

ductile cast iron crankshafts. The S-N predictions proved to be a more accurate life

prediction method, providing reasonable results for both the forged steel and cast

iron crankshafts. The strain-life predictions also provided reasonably accurate

estimations for the fatigue life of the forged steel crankshafts and less accurate,

however conservative, estimations for the ductile cast iron crankshafts.

14. The accuracy of fatigue life predictions using the S-N or the strain-life approach is

strongly influenced by an accurate estimation of notch sensitivity of a material.

Using a low notch sensitivity for the ductile cast iron crankshaft (q = 0.2) as

suggested in the literature resulted in life predictions that did not agree with the

experimental data. When low notch sensitivity was assumed the predictions

overestimated the results while high notch sensitivity underestimated the results.

146

REFERENCES

Asi, O., 2006, “Fatigue Analysis of a Crankshaft Made from Ductile Cast Iron,” Fatigue Failure Analysis, Vol. 13, pp. 1260-1267 ASTM Standard E8-04, 2004, “Standard Test Methods for Tension Testing of Metallic Materials,” Annual Book of ASTM Standards, Vol. 03.01, West Conshohocken, PA, USA. ASTM Standard E23-02a, 2004, “Standard Test Methods for Notched Bar Impact Testing of Metallic Materials,” Annual Book of ASTM Standards, Vol. 03.01, West Conshohocken, PA, USA. ASTM Standard E83-02, 2004, “Standard Practice for Verification and Classification of Extensometer System,” Annual Book of ASTM Standards, Vol. 03.01, West Conshohocken, PA, USA. ASTM Standard E606-92, 2004, “Standard Practice for Strain-Controlled Fatigue Testing,” Annual Book of ASTM Standards, Vol. 03.01, 2004, West Conshohocken, PA, USA. ASTM Standard E646-00, 2004, “Standard Test Method for Tensile Strain-Hardening Exponents (n-Values) of Metallic Sheet Materials,” Annual Book of ASTM Standards, Vol. 03.01, West Conshohocken, PA, USA. ASTM Standard E739-91, 2004, “Standard Practice for Statistical Analysis of Linear or Linearized Stress-Life (S-N) and Stain-Life (ε-N) Fatigue Data,” Annual Book of ASTM Standards, Vol. 03.01, West Conshohocken, PA, USA. Bayrakçeken, H., Tasgetiren, and S., Aksoy, F., 2006, “Failures of Single Cylinder Diesel Engines Crankshafts,” Fatigue Failure Analysis, Vol. 14, pp. 725-730. Behrens, B. A., Reinsch, S., Specker, A., and Telkamp, K., 2005, “Further Development in the Precision Forging Technology for High Duty Automotive Parts,” In MPMD Sixth Global Innovations Proceedings Trends in Materials and Manufacturing Technologies for Transportation Industry and Powder Metallurgy Research and Development in the Transportation Industry, San Francisco, CA, USA, The Minerals, Metals, and Materials Society. Chatterley, T.C. and Murrell, P., 1998, “ADI Crankshafts - An Appraisal of Their Production Potentials,” SAE Technical Paper No. 980686, Society of Automotive Engineers, Warrendale, PA, USA.

147

Chien, W.Y., Pan, J., Close, D., and Ho, S., 2005, “Fatigue Analysis of Crankshaft Sections Under Bending with Consideration of Residual Stresses,” International Journal of Fatigue, Vol. 27, pp. 1-19. Damir, A.N., Elkhatib, A., and Nassef, G., 2007, “Prediction of Fatigue Life Using Modal Analysis for Grey and Ductile Cast Iron,” International Journal of Fatigue, Vol. 29, pp. 499-507. Fergusen, C. R., 1986, “Internal Combustion Engines, Applied Thermo Science,” John Wiley and Sons, New York, NY, USA. Grum, J., 2003, “Analysis of Residual Stresses in Main Crankshaft Bearings after Induction Surface Hardening and Finish Grinding,” Journal of Automobile Engineering, Vol. 217, pp. 173-182. Heitmann, W.E., August 14, 2006, Private Communication. http://www.tpub.com/engine3/en3-53.htm http://en.wikipedia.org/wiki/Four-stroke_cycle Jensen, E.J., 1970, “Crankshaft Strength Through Laboratory Testing,” SAE Technical Paper No. 700526, Society of Automotive Engineers, Warrendale, PA, USA. Juvinall, R.C. and Marshek, K.M., 1991, “Fundamentals of Machine Design Components,” 2nd Edition, John Wiley & Sons, New York, NY, USA. Laus, L. and Heitmann, W.E., February 15, 2007, Private Communication. Montazersadgh, F.H., 2007, “Stress Analysis and Optimization of Crankshafts Subjected to Dynamic Loading,” Masters Thesis, The University of Toledo, Toledo, OH, USA. Montazersadgh, F. and Fatemi, A., 2007, “Dynamic Load and Stress Analysis of a Crankshaft,” SAE Technical Paper No. 2007-01-0258, Society of Automotive Engineers, Warrendale, PA, USA. Park, H., Ko, Y. S., and Jung, S. C., 2001, “Fatigue Life Analysis of Crankshaft at Various Surface Treatments,” SAE Technical Paper No. 2001-01-3374, Society of Automotive Engineers, Warrendale, PA, USA. Pichard, C., Tomme, C., and Rezel, D., 1993 “Alternative Materials for the Manufacture of Automobile Components: Example of Industrial Development of a Microalloyed Engineering Steel for the Production of Forged Crankshafts,” In Proceedings of the 26th ISATA International Symposium on Automotive Technology and Automation, Aachen, Germany. Shigley, J.E., and Mitschke, C.R., 2002, “Mechanical Engineering Design,” 5th Edition, McGraw-Hill, Inc., Boston, MA, USA. Silva, F.S., 2003, “An Investigation into the Mechanism of a Crankshaft Failure,” Key Engineering Materials, Vols. 245-246, pp. 351-358, Trans Tech Publications, Switzerland.

148

Spiteri, P., Ho, S., and Lee, Y., 2007, “Assessment of a Bending Fatigue Limit for Crankshaft Sections with Inclusion of Residual Stresses,” International Journal of Fatigue, Vol. 29, pp. 318-329. Stephens, R.I., Fatemi, A., Stephens, R.R., and Fuchs, H.O., 2000, “Metal Fatigue in Engineering,” 2nd Edition, John Wiley & Sons, Inc., New York, NY,USA. Wang, Z., Xu, J., Bao, G., Zhang, J., Fang, X., and Liu, T., 2007, “Fabrication of High-Powered Diesel Engine Crankshafts by Electro-slag Casting Process,” Journal of Materials Processing Technology, Vol. 182, pp. 588-592. Williams, J. and Fatemi, A., 2007, “Fatigue Performance of Forged Steel and Ductile Cast Iron Crankshafts”, SAE Technical Paper No. 2007-01-1001, Society of Automotive Engineers, Warrendale, PA, USA. Williams, J., Montazersadgh, F.H., and Fatemi, A., 2007, “Fatigue Performance Comparison and Life Prediction of Forged Steel and Ductile Cast Iron Crankshafts,” 27th Forging Industry Technical Conference, Fort Worth, TX. Zoroufi, M. and Fatemi, A, 2005, “A Literature Review of Durability Evaluation of Crankshafts Including Comparisons of Competing Manufacturing Processes and Cost Analysis,” 26th Forging Industry Technical Conference, Chicago, IL.

All rights reserved 2008 by David Sparkman, MeltLab Systems

Ductile Iron Microstructure by Thermal Analysis

One promising new technology is the rapid classification of microstructure by thermo

analysis. Current technology consists of either pouring a sample with two ears on it for

micro analysis or creating a micro lug in the runner system of a mold and digging the lug

out for grinding and polishing. Both techniques are “after the fact” or as we often say,

“post mortem” since a bad result results in scrapping already poured castings.

For a while there has been an analysis method by NovaCast, and now cloned by Elkem to

analysis a treated ductile iron sample for the propensity toward shrinkage. But there is so

much more that can be learned though thermo analysis.

NovaCast identified several characteristics of the thermal curve as pertaining to good

castings. First they wanted to see a quick rise in the eutectic temperature as compared to

the length of the eutectic arrest (T2 to T3 Ratio). This depended on good nodule count

and is soundly based in metallurgy.

Second, they looked for a sharp drop in temperature at the end of the eutectic. This shows

up as a strong dip in the first derivative, or a sharp peak in the cooling rate. This arrest is

due to a large stress in the grain boundaries. The lack of this stress means that shrinkage

has relieved the stresses. The stresses may affect fatigue life, but no one has investigated

that aspect. For sure, shrinkage will also greatly reduce fatigue life.

With some previously published information from MeltLab introducing the concept of

the slope of the temperature fall off after eutectic as an indicator of grain size, Elkem has

adopted reporting that slope.

Additional features in Thermal Analysis

The first concept is that not all ductile foundries need the same kind of ductile iron to

make quality castings. The treatments by NovaCast and Elkem are aimed at the average

foundry making thinner section castings and do not discriminate between hypo-eutectic,

eutectic, and hyper-eutectic irons. The published examples all appear as eutectic curves,

though closer examination is needed to pick out the near eutectic curves.

Let’s start with some definitions of terminology. The term “eutectic” means both lowest

melting point, and single arrest. Many may assume that an eutectic curve (single arrest)

corresponds with the Iron Carbon phase diagram C.E. of 4.33. This is not always so for

two reasons: first, magnesium suppresses the formation of graphite which delays the

formation of graphite and allows an austenite liquidus to form higher than 4.33. Second,

the instrumentation used by both NovaCast and by Elkem cannot always find an

austenitic arrest hidden in a eutectic arrest. Nor do they spot the small graphitic liquidus

arrests of extremely hyper-eutectic iron.

All rights reserved 2008 by David Sparkman, MeltLab Systems

Here is a typical “eutectic curve” with what appears to be a single arrest. The lack of

resolution hides the austenite liquidus present.

By showing the rate of cooling curve below, the austenite liquidus is seen before the

general eutectic arrest.

Actually it is a hyper-eutectic at 4.49C.E.

but with an austenitic liquidus due to magnesium

The above curve without and with cooling rate demonstrates the need for a close look at

the cooling rate to pick up the small but significant austenitic arrest. Nodularity is not as

good as can be had as can be seen by the fine roughness in the green curve during

eutectic, and there is an indication of shrinkage at 1:50 minutes into the sample. More on

this later.

All rights reserved 2008 by David Sparkman, MeltLab Systems

The following curve is a true eutectic curve, even though the C.E. is 4.43%.

Single arrest Ductile Iron with 4.43 C.E.

This kind of iron is idea for normal section castings and provides good shrink control for

reasons to be explained later. Nodularity is quite good in this sample as can be seen by

the smoothness of the green curve.

The final type of curve is a hyper-eutectic arrest with a graphitic liquidus, and often with

an austenitic liquidus as well.

Hyper-eutectic iron with graphitic arrest at 4.60% C.E. and other problems

All rights reserved 2008 by David Sparkman, MeltLab Systems

The graphitic arrest is quite weak, but shows up as a clear arrest in the cooling rate.

Looking only at the red temperature curve, it would be easy to miss. But look at the fall

out. We have a good smooth green curve during the eutectic indicting good nodule shape,

but we have two very bad arrests between eutectic and end of freezing. The first arrest at

2:30 minutes into the analysis is exothermic or heat producing. At this point in the curve,

that indicates a D or E flake formation commonly referred to as carbides. The second

smaller arrest is endothermic or heat adsorbing. This is actually shrinkage forming in the

sample.

Hyper-eutectic iron with graphitic arrest at 4.60% C.E. showing carbide arrest and shrink

All rights reserved 2008 by David Sparkman, MeltLab Systems

Shrink in ductile iron thermo analysis

In steel, the shrinkage rule is 10%. This means that the pattern will be designed for all

dimensions to be 10% over stated so that when the casting cools, it will have shrunk to

near the right size. In ductile iron, graphite occupies between 9 and 11% of the volume,

while in liquid ductile iron, carbon occupies no volume. The miracle of ductile iron is

then that graphite grows in the solidifying material to almost perfectly match the

shrinkage. The key word is “almost’. If we get enough graphite to come out of the liquid

at the right time, we produce sound castings. If not, we find shrinkage in our castings.

The purpose of the gating system is to feed liquid into the cavity as the liquid cools.

Liquid iron at 2200 degrees F has less volume per pound than liquid iron at 2300 degrees

F. We want the gates to freeze off before the graphite starts to grow in the mold cavity.

This traps the increase in volume inside the casting; instead of pushing iron back out into

the runner system and making solid risers. This works best if no significant amount of

graphite comes out in the liquid. Obviously if you do have graphite forming in the liquid

you would expect to find a bimodal distribution in your nodule sizes. There is some

discussion if the larger graphite can also grow in the early austenite arrest and so give rise

to bimodal distribution.

The effect of Pearlite on Shrinkage and the thermal analysis

Second, some carbon does remain in the iron after solidification. Some is in the obvious

form of carbides or pearlite, but some remains in solution below microscopic detection.

This is proven by the effect of heat treating causing an increase in graphite volume even

in fully ferritic irons. Increased nodule counts contribute strongly to more complete

graphitization. And obviously carbide stabilizers decrease graphitization. In the thermo

analysis curve, the prolonged eutectic plateau is a good indicator of the degree of

graphitization, and its counter: an early end of the plateau, a good indication of pearlite

content.

All rights reserved 2008 by David Sparkman, MeltLab Systems

Here are two cooling rates, the top is ferritic iron, and the bottom is pearlitic iron. The

energy production of the top curve is significantly greater toward the end of the eutectic

arrest.

In this example of an 80% pearlite, the End of Graphitization is lifted up to about 1.2

degrees per second rate of cooling. Not near as much graphite came out of this iron.

Please remember that the carbon is still all in the matrix, the pearlite does not form until

1400-1450 F. But because it is still in the matrix as a form of austenite, it’s energy has

not yet been released.

The above two examples show the temperature curve of a pearlitic curve (top) and a

pearlitic curve (bottom).

All rights reserved 2008 by David Sparkman, MeltLab Systems

This is an example of Ferrite ductile (<10% pearlite). The 4th derivative is shown in black

so you can see the bend associated with the end of graphitization. This point is at about

0.4 degrees per second rate of cooling, so the material has been able to maintain a much

flatter eutectic due to higher energy production. Of course the ferritic iron also has higher

nodule count, you don’t want to over inoculate pearlitic iron. It gets soft.

All rights reserved 2008 by David Sparkman, MeltLab Systems

In this ductile sample, I would like to introduce the Steady State Cooling Rate concept.

When no crystallizations or solid state transformations are occurring, then the metal

looses heat at a steady, mathematically describable rate due most exclusively to radiation

and conduction. Here I have drawn straight lines to represent the SSCR of the iron liquid

(SSCRL), and the iron solid (SSCRS). The lines are slightly offset so that you can see

how they parallel the actual readings. In reality the lines show a slight bend and are

logarithmic. The reason for the different slopes is that conductivity and radiation rates

change with structure.

The usefulness of this concept is that these lines can serve as a boundary or delimiter for

energy production. Combined together, they define what in the literature is called the

zero line. Here is an example of how those lines extend to make energy boundaries.

All rights reserved 2008 by David Sparkman, MeltLab Systems

The start of Liquidus is where the curve breaks away from the SSCRL cooling rate and

begins to produce heat. Linking this with the SSCRS gives us the upper boundary for

calculating the energy of solidification, as well as a lower boundary to calculate the

energy adsorbed by the grain boundary solidification. These are very important

measurements for calculating the quality of the ductile iron. Knowing the total energy of

solidification allows us to ratio any arrest such as carbides or shrinkage against the total

energy, and come up with a percentage of energy for each type of arrest. While the

energy for each type of reaction may differ (different specific heats), knowing that one

sample had carbides of 1% of the total energy vs. one with 0.1% of the total energy

allows us to understand the severity of the problem.

Understanding where the lines come from, let’s look at the meaning of the area under the

Solidus point known as the grain boundary stress area (GBSA). This is an endothermic

reaction that sucks up energy due to disorder (stress) in the grain boundaries between

crystals. Sometimes it is quite large, sometimes it is quite small. Nova Cast has used the

height of the point as a shrink indicator based on their studies, and for good reasons.

There should always be stress in the grain boundaries after solidification. In the long

term, heat treating removes these stresses by increasing the amount of graphite present.

But in the short term, there is too little iron for too much volume and there is only one

way to reduce the stress – shrinkage. The shrinkage may take the form of a suck-in on the

top or sides of the thermal analysis cup, or by forming actual shrinkage inside the casting.

The shrinkage shows up on the thermal analysis curve, the suck-in does not.

So a quick estimate of this stress is the height of the Solidus point. But since that can vary

with how well the cup is filled, a better estimate can be the area between the Rate of

Cooling curve and the zero line ratio against the total energy of solidification.

All rights reserved 2008 by David Sparkman, MeltLab Systems

In summary, shrinkage has many causes, graphitic liquids’ can remove significant

amounts of carbon, the carbon/C.E. content may be too low, inoculation can alter when

and how much of the graphite comes out. Pearlitic irons are designed to retain from 0.3 to

0.6% carbon in the matrix which later comes out as between 40 and 80% pearlite. (40%

times 0.8 Carbon in 100% pearlite equals 0.32% retained carbon.)

All rights reserved 2008 by David Sparkman, MeltLab Systems

Statement of thanks (With a slight nod to Bill O’Reilly for an attitude check)

It is customary in research papers to list a biography of all references searched. While this

is a good practice and recommended and even required by those who pass judgment on a

paper’s worthiness, I find it difficult to recommend most research on this subject as they

contain so many errors and misconceptions. While a good grounding in metallurgy is

necessary to understand some of these concepts, and history is an interesting field, the

study of disease and astronomy before the age of the invention of the optical lens by

Galileo produced mostly useless information.

In metallurgical thermo analysis, the lack of clear and sharp derivatives made analysis

only guesswork. A paper by a studied researcher mentioned the elusive graphite arrest. It

is visible in DTA but not in normal thermo analysis. Therefore many researchers do not

believe that you can produce graphite in a liquid in normal foundry iron. But you can see

it with the modern tools of MeltLab.

My Heroes My thanks go out to Jeff Burk, Dick Heine and Bill Shaw for first interesting me in the

possibilities of TA, to Carl Loper for introducing me to grain boundary freezing, to

Toban xxx for an interesting discussion on graphitization of ferritic and pearlitic irons.

And finally to Doru Stephenescu for discussions on the Zero curve.

Now for the Pinhead section My undying gratitude goes to ElectroNite for copying MeltLab’s Solidus Point and for

their continuing hatred and putdowns of MeltLab that have only encouraged me to

embarrass them further. Last overheard from them: “what are we going to do about those

people in Virginia?”

My undying gratitude goes to greedy Minco and OCC for selling our products to a few

customers while secretly having programmers attempt to duplicate MeltLab. Sorry those

products failed.

And finally my fond farewell to NovaCast who through kickbacks (up to 20% according

to one Metallurgist) and other financial incentives, tried to sell an overpriced box (40k+).

They exchanged heavy advertising in the DIS for preferred treatment in making multiple

presentations that included nothing but more advertising. They also succeeded in hiring

away a technical writer/researcher from us to get at MeltLab technology. Their attempt,

with investor money, to control the market for ductile thermo analysis is now coming to

an end. Their promises to investors were just smoke and mirrors. Now that Elkem Metals

(a good company with value in its products) has produced an ATAS Light product (there

is a lawsuit over that), and with the introduction of MeltLab for Ductile, the overpriced

ATAS product will soon, if not already, be history.

Journal of Minerals & Materials Characterization & Engineering, Vol. 7, No.4, pp 307-316, 2008

jmmce.org Printed in the USA. All rights reserved

307

To Study the Effect of Austempering Temperature on Fracture Behaviour of

Ni-Mo Austempered Ductile Iron (ADI)

Vikas Chawla1*, Uma Batra

2, D. Puri

3, Amita Chawla

4

1Mechanical Engineering Department, L.L.R.I.E.T, Moga, Punjab, India

2Metallurgy Department, P.E.C, Chandigarh, India

3Metallurgical & Materials Engineering, I.I.T. Roorkee, India

4Chemistry Department, Govt. Brijindra College. Faridkot, Punjab, India

*Corresponding author: E-mail: edwalesir@rediff.com,

Phone: +91-9719749154, Fax: +91-1332-285243

ABSTRACT

Austempered Ductile Iron (ADI) can be as twice as strong as standard spheroidal iron at the

same level of toughness. It responds to work-hardening surface treatments and exhibits excellent

fatigue and wear property. There is extensive work done on the fracture of steel with ferrite

or/and austenite structure, but little on fracture behaviour of ADI whose microstructure also

comprises austenite and ferrite but with graphite nodules in the matrix. The present work is

aimed in this direction. The fracture behavior of Ni-Mo ADI is studied. It is found that the crack

always originates from graphite nodules and the matrix affects the propagation path.

Keywords: Austempered ductile iron (ADI), Fractography, Dimpled structure, Fracture,

Fracture mode.

1. INTRODUCTION

Austempered Ductile Iron (ADI) has ‘come of age’ during its brief history. ADI production is

expected to grow at an annual rate of at least 5% [1]. The microstructure of ADI also comprises

austenite and ferrite (in the form of bainitic ferrite), but with graphite nodules in the matrix. The

market of ADI is extremely large. Their attractive properties make them desirable not only for

the manufacture of existing components with improved performance but also for competing with

other materials in new applications. Advantages of ADI include high strength, ductility, wear

resistance, toughness, better machinability, high damping capacity and reduced weight in

comparison with forge steel. ADI has been widely used for engineering components such as

308 Vikas Chawla, Uma Batra, D. Puri, Amita Chawla Vol.7, No.4

gears, crankshaft, vehicle components, sprockets, and cutting tools. The matrix of ADI can

withstand a certain amount of deformation before fracture during tensile and impact testing.

However, the graphite nodules in the matrix cannot deform and hence are barriers to matrix

deformation, which give rise to crack initiation. The crack propagation and the fracture mode of

ADI are influenced by the orientation of bainitic ferrite needles with respect to the load direction

and also by the presence of carbide particles inside the needles or at interfaces.

2. EXPERIMENTAL PROCEDURE

The tensile and impact test specimens of standard dimensions (as per ASTM) were machined out

of Ni-Mo ductile iron casting (in the shape of cylindrical bar). The chemical composition (wt %)

is shown in Table 1.

Table 1. Nominal chemical composition (wt %) of Ni-Mo Ductile Iron casting.

C Si Ni Mo Mn Mg P S Fe

3.43 3.02 1.16 0.43 0.21 0.12 0.016 0.007 balance

Subsequently these specimens were annealed at 720°C for 4 hrs to homogenize the structure and

to achieve uniform distribution of alloying elements. Figure 1 shows the microstructure of Ni-

Mo ductile iron after annealing.

Figure 1. Ni-Mo Ductile Iron annealed at 720°C for 4 hrs

After annealing, all the specimens were austenitized at 900°C for 1hr and then austempered at

different temperatures in a salt bath. Table 2 indicates the designation of the specimens as per the

austempering temperature. The composition (wt %) of salt bath is shown in Table 3.

Table 2. Designation of specimen under study.

Specimen Designation

Austempered at 270°C for 1 hr. A-1

Austempered at 330°C for 1 hr. A-2

Austempered at 380°C for 1 hr. A-3

Vol.7, No.4 Study the Effect of Austempering Temperature 309

Table 3. Composition (wt %) of salt bath.

KNO3 NaNO2 NaNO3

40 55 05

All the specimens austempered at different temperatures were then fractured under tensile and

impact loading. The fractured pieces were stored in plastic bags before fractography, in order to

prevent any chemical or mechanical damage to them.

The fractured surfaces were analyzed by Scanning Electron Microscope (SEM) to obtain

fractographs at various locations. For fractomicrography, the specimens were sectioned in the

direction perpendicular to the fractured surface, and then the specimens were polished.

Subsequently the samples were etched in Nital (97ml CH3OH, 3ml nitric acid). After polishing

and etching the specimens were observed under SEM and fractomicrographs were taken in order

to view the crack propagation path.

3. OBSERVATIONS

3.1 Fractography

Visual, with Optical Microscope, SEM and TEM observations of as fractured surfaces, is termed

as Fractography. The effect of microstructure has been observed on the fractured surface

appearance i.e. fractography. Figure 2 and 3 show the fractures surface appearance for the

specimen A-1 broken in impact and tensile test respectively. The dimpled surface is observed.

These dimples indicate that the fracture may have occurred by the phenomenon called void

coalescence i.e. separation of the material internally, forming voids which then join to develop

the fracture surface. The shallow dimples can be seen at various locations, which indicate high

strength & low ductility of the material [2, 3].

The fractographs of specimen A-2, broken in impact test and tensile test are shown in Figures 4

and 5 respectively. These fractographs shows the dimples at various locations which indicate the

fracture may have occurred by void coalescence. The dimples are deep as compared to A-1

specimen, which indicate low strength and more ductility [2, 3].

The fractographs of specimen A-3, broken in impact test and tensile test are shown in Figures 6

and 7 respectively. These fractographs shows signs of cleavage as well as of void coalescence.

The fractographs show river like pattern as well as dimples at various locations. This indicate

that the fracture my have occurred by the mixed phenomenon i.e. quasi cleavage fracture

mechanism [4]. According to which, the fracture may have occurred by cleavage at some foreign

particle subsequently separated from the matrix by void coalescence.

310 Vikas Chawla, Uma Batra, D. Puri, Amita Chawla Vol.7, No.4

3.2 Fractomicrographv

Fractography gives information about the nature of fracture whereas the fractomicrography is the

study of the surface, which is being cut perpendicular to fractured surface. After sectioning and

mounting, the fractured specimens were analyzed by SEM in order to determine the crack

initiation and propagation, graphite nodule shape, size, distribution and matrix structure.

Fractomicrographs of specimens A-1, A-2 and A-3 are shown in Figures 10, 11 and 12

respectively.

3.3 Matrix Structure

The matrix of ADI is a complex mixture of bainitic ferrite and austenite, where austenite is the

basic phase. On austempering at 250 - 330°C, the matrix of ADI comprises the banitic-ferrite

needles with carbide particles inside them and the rest is austenitic. Whereas austempering at 330

- 450°C, the matrix of ADI observed to consist bainitic-ferrite needles without carbide particles

and carbon rich / stabilized austenite, as carbon diffuses to austenite the basic phase [5].

Several bainitic ferrite needles or plate- lets have the same orientation, forming a cluster of

bainitic ferrite needles or platelets. Each cluster has a particular orientation, as shown in Figure

8.

Under uniaxial external load the orientation relationship between bainitic ferrite needles and the

applied load direction can be classified into three types: -

(i) The longitudinal direction of cluster of bainitic ferrite needles is parallel to the loading

direction 'P', as shown in Figure 9(a).

(ii) The longitudinal direction of a cluster of bainitic ferrite needles is perpendicular to the

loading direction 'P', as shown in Figure 9(b).

(iii) The longitudinal direction of a cluster of bainitic ferrite makes angle 'θ' with the loading

direction 'P', as shown in Figure 9(c).

Most of the clusters of bainitic ferrite needles belong to category as shown in Figure 9(c), but

with different angles.

Figure 2: Fractographs of specimen austempered Figure 3: Fractographs of specimen austempered

at 270°C for 1 hr. and broken in Impact test. at 270°C for 1 hr. and broken in Tensile test.

Vol.7, No.4 Study the Effect of Austempering Temperature 311

Figure 4: Fractographs of specimen austempered Figure 5: Fractographs of specimen austempered

at 330°C for 1 hr. and broken in Impact test. at 330°C for 1 hr. and broken in Tensile test.

Figure 6: Fractographs of specimen austempered Figure 7: Fractographs of specimen austempered

at 380°C for 1 hr. and broken in Impact test. at 380°C for 1 hr. and broken in Tensile test.

Figure 8: The orientation of clusters of bainite Figure 9: Systematic simplified diagram showing the

ferrite needles with applied load direction “P.” possible orientation relationship between bainitic

ferrite needles with applied load direction “P.”

-->P P<- -

312 Vikas Chawla, Uma Batra, D. Puri, Amita Chawla Vol.7, No.4

Crack propagation from a graphite

nodules

Crack propagation along the

ferrite/austenite interface

Crack propagation by cutting

through bainitic ferrite needles

Crack propagation along the

interfaces and cutting through

banitic ferrite needles

Figure 10: Fractomicrograph showing

crack propagation path in the specimen

having lower bainite microstructure (i.e.

austempered at 270°C for 1hr). The

direction of “P” indicate the load direction

→ P

→ P

→ P

→ P

→ P

→ P ←

←

←

←

←

←

Vol.7, No.4 Study the Effect of Austempering Temperature 313

Figure 11:Fractomicrograph showing crack Figure 12: Fractomicrograph showing crack

Initiation & propagation in the specimen initiation at second phase paricle and

austempered at 330°C for 1 hr. propagation in the specimen austempered at

380°C for 1 hr.

4. DISCUSSION

Graphite nodules being discontinuities in the ADI matrix, give rise to much higher stresses

around the graphite nodules during elastic deformation [6], and hence causes crack initiation at

graphite / matrix interface.

For a structure similar to Figure 9 (b), the bainitic ferrite, austenite and the interfaces of

ferrite/austenite will undergo similar external tensile stress. Atomic mismatch at ferrite/austenite

interface decreases the tensile stress bearing capacity as compared to bainitic ferrite and austenite

individually. As a result, the cracks, which originate from the graphite nodule, usually propagate

along the interfaces of ferrite/austenite as indicated in Figure l0. For a structure similar to Figure

9(a) the bainitic ferrite and austenite will deform with load. As austenite has better ductility than

banitic ferrite needles, thus can sustain more deformation [1]. So banitic ferrite needles will

break first. The austenite deforms even after the fracture of the bainitic ferrite needles. This crack

will proceed by cutting the bainitic ferrite needles.

For the most common bainitic ferrite needles structure, the crack take the easiest way to

propagate, as is observed in Figure l0. The crack propagates along the interfaces of bainitic

ferrite/austenite when angle between the applied load direction & longitudinal direction of

cluster is greater than 45° and sometimes cut through the needles when angle is less then 45°.

For microstructure corresponding to the specimen A-1, which is lower bainite, carbide particles

appear in the bainitic ferrite needles, which act as barrier to ferrite slip. The carbide is hardest

and brittle as compared to ferrite and austenite in the matrix of ADI. There is high stress

concentration around carbide particles, during deformation under load. So the bainitic ferrite

314 V. Chawla, U.Batra, D. Puri and A. Chawla Vol.7, No.4

needles having carbide particles inside them facilitate the crack to pass through bainitic ferrite

needles or platelets [1], as observed in Figure 10. This creates opportunities for cracks to select

an easy way to propagate. Also the carbide particles deflect the path of crack propagation. This

results in the shallow dimpled fractured surface (as observed in figures: 2 & 3), which is due to

high strength of the material [3].

For microstructure corresponding to the specimen A-2, crack initiation and propagation is shown

in Figure 11. The microstructure consists of lower bainite (which consist of bainitic ferrite

needles with carbide particles) and retained austenite. The dimples appeared deeper then A-1,

which indicate this material is more ductile. The crack is initiated at graphite/matrix interface

and propagates along bainitic ferrite needles/austenite interfaces or cut through ferrite needles,

depending upon the orientation relationship of bainitic ferrite needles with applied load direction.

For the microstructure corresponding to specimen A-3, the ADI consist of bainitic ferrite needles

without carbide and high carbon stabilized austenite i.e. upper bainite is the microstructure.

Upper bainite has lower strength then lower bainite microstructure, as ferrite is the softest

structure. Figures 6 and 7 (showing the signs of cleavage i.e. river like pattern as well of void

coalescences with dimpled structure at some locations) give indication of Quasi-cleavage

fracture in the specimen A-3, which is supported by the Figure 12. This shows that the crack is

initiated at second phase particle, which is hard & brittle as compared to the matrix of ADI

(which is soft & comparatively ductile). So brittle & cleavage fracture has occurred on the

second phase particle, then separation of connecting material by void coalescence. Figure 12

show the path of propagation of crack.

According to observations the cracks are originating from graphite nodules in ADI. The easiest

propagation paths for cracks are the interfaces between ferrite and austenite, because of atomic

mismatch at the ferrite/austenite interfaces. The orientation of a cluster of bainitic ferrite needles

and the presence of precipitated carbide in the matrix can influence the crack path. However, the

orientation of the longitudinal direction of banitic ferrite needles is random and does not

influence the fracture mode of ADI. Precipitated carbides in the matrix of ADI do not

significantly influence the fracture characteristic of ADI. In order to analyze the crack

propagation path, Fan et al. [6] have explained the crack propagation with the help of a model as

shown in Fig.13. The same model is verified in our study, in which two clusters of banitic ferrite

needles between two graphite nodules, one nearly parallel to applied load direction and the other

nearly perpendicular to the applied load direction can be assumed. The effect of precipitated

carbide on the crack propagation path and the fracture mode of ADI can be explained with a

model as shown in Figure 13 [6].

Vol.7, No.4 To Study the Effect of Austempering Temperature 315

Figure 13. Two models of crack propagation in ADI [5].

Model shows:

Micro voids at graphite -matrix interfaces: -1

Bainitic ferrite needles: -2

Carbide particles in bainitic ferrite needles: -3

Possible crack path: -4

Model (I) in Figure 13 shows no carbide appear in the ADI matrix crack often pass along the

ferrite/austenite interfaces for which the needles have greater angle than 45° with the applied

load direction. However, if the longitudinal direction of the cluster of bainitic ferrite needles

tends to be parallel to the loading direction, cracks may cut through the ferrite needles (as ferrite

is softest microstructure) & austenite. In this case the fracture mode should be ductile, but due to

presence of second phase particles the fracture mode is Quasicleavage as explained earlier.

Model (II) in Figure 13 show carbide in the bainitic ferrite needles or platelets, which act as

barriers to ferrite slip, results in higher stress around carbide particles [6]. Large number of stress

concentration locations in the needles creates particles, which further creates more opportunities

for crack to pass through needles or platelets. As the carbide particles are harder so undergoes

fracture without deformation and also deflect the crack path. This creates opportunity for crack

to select an easy way to propagate. This result has further been supported as the fractography

shows fracture in ductile mode.

Figure 13 gives just one of the main possible propagation paths of the crack. Although we cannot

predict the particular propagation path, the observed results can help us understand and develop

the appropriate microstructure of ADI.

CONCLUSIONS

The following conclusions can be drawn from the above study

1. Crack always originates from graphite nodules in ADI.

2. The easiest path of propagation of a crack is along the austenite/ ferrite interfaces due to

atomic mismatch.

316 V. Chawla, U.Batra, D. Puri and A. Chawla Vol.7, No.4

3. The longitudinal direction of bainitic ferrite needles can be parallel, perpendicular, or inclined

at angle 'θ' with the load direction.

4. The propagation path of a crack in ADI depends upon the orientation relationship of bainitic

ferrite needles with the applied load direction.

5. The fracture mode in lower bainite structure is ductile mode and the fractographs shows

shallow dimpled structure, which indicates the high strength of lower bainite microstructure of

ADI.

6. The fracture mode in upper bainite microstructure is Quassi cleavage or mixed mechanism i.e.

the fracture firstly occurs by cleavage then separation of material by void coalescence. The

fracture of second phase particle occurs by cleavage due to lack of ductility and of the matrix by

void coalescence (plastic deformation).

7. Carbide particles in bainitic ferrite needles promote the passage of cracks through the ferrite

needle, but do not significantly influence the fracture mode of ADI.

REFERENCES

[1] R.J Warrick “Application of Ductile Iron Castings”. ASM Technical reports system no. 76-

45.

[2] Calangels and Heiser: Metallurgical Failure: A Wiley-Interscience publications.

[3] J.C Morrison, “What’s’ in a name-nickel and ductile iron”. Technical paper in Indian

Foundry Journal, 1998, Vol. 44(12).

[4] Ashok Chaudhary and Charlie R. Brooks “Metallurgical Failure Analysis,” McGraw Hill

publisher.

[5] Ray Elliot: Cast Iron Technology, Jaico publishing house.

[6] Z.Fan and R.E Smallman (1994) “Some observations of Austempered Ductile Iron”. Scripta

Metallurgica et Materials journal, Vol.31 (2), 1994

1

Developments in Cast Iron

Metallurgical Treatments

by

Dr. Torbjørn Skaland

Elkem ASA, Research

Kristiansand, Norway

2

Abstract

The objective of the present paper is to review some selected

theoretical and practical implications in metallurgical

treatment and solidification of cast iron. Treatment of

liquid iron with ferrosilicon based alloys have been given

specific focus, and plausible explanations to the mechanisms

of graphite spheroidization and inoculation are given.

Effects of nodularizing (magnesium treatment) on nucleation

potentials and iron quality in general are discussed. A

comparison of different commercial treatment methods and

nodularizing agents (metals and alloys) are shown to reflect

the importance of magnesium treatment on the subsequent

inoculation effectiveness and resulting solidification

characteristics. An optimized magnesium treatment process is

described which gives good process economy, consistency,

environment and final casting quality.

Research on the inoculation mechanisms in ductile cast iron,

with particular emphasis on the aspects of heterogeneous

nucleation of graphite at inclusions are reviewed. It is

shown that the majority of inclusions in ductile cast iron

are primary or secondary products of the magnesium treatment

(i.e. sulphides and oxides). After inoculation, special

hexagonal silicate phases form at the surface of oxide

inclusion. The presence of these phases will enhance the

nucleation potency of the inclusions with respect to

graphite. The theoretical analysis of reactions during both

magnesium treatment and inoculation is in close agreement

with experimental and practical observations.

3

1. Introduction

One of the most important stages in the iron founding process

is the economic production of liquid iron and its

metallurgical treatments in preparation for pouring into the

mould. This involves maintaining compositional and

temperature control over the liquid during melting in order

to achieve the correct composition for the specified iron,

the correct graphitization potential, and the correct state

of nodularizing and inoculation to ensure a sound casting of

the desired structure and the required properties.

Many of the different elements added to iron is supplied

through specialty ferroalloys as nodularizers and inoculants.

Process developments and product quality shifts in cast iron

metallurgy rely much on improved compositional control. This

again is reflected on ferroalloys in terms of stricter

quality specifications for greater uniformity and higher

purity. The present and future role of ferroalloys has to be

viewed on the background of the needs of the iron founding

industry since further requests for alloy modifications or

possible new grades clearly derive from this industry.

Magnesium is the most popular nodularizing agent, and it is

usually added in multi-component alloy form. These components

contain additions to reduce reaction violence, promote

graphite spheroidization, neutralize the effect of impurities

on graphite morphology, and control the matrix structure. The

most common alloys for nodularizing iron are ferrosilicon

containing 3 to 12 per cent magnesium.

Inoculation is a means of controlling the structure and

properties of cast iron by minimizing undercooling and

4

increasing the number of nucleation events during

solidification. The most effective inoculants today are

ferrosilicon alloys containing small quantities of elements

such as Ca, Al, Ba, Sr, Zr, and Ce.

The present paper is an attempt to outline basic

understanding and practical findings for optimizing ductile

iron production and properties. Nodularizing and inoculation

treatment will be covered separately, although the important

connection between the two treatment actions will also by

described.

2. Basic Effects of Nodularizing Treatment

Nodularizing, or magnesium treatment, of cast iron is a means

of modifying the solidification structure so that the

graphite phase precipitates and grows as spherical particles

instead of flakes, thus resulting in a cast iron with

significantly improved mechanical properties. The

nodularizing action can be regarded as a simultaneous

desulphurizing and deoxidizing treatment, where elements

having strong affinity to both sulphur and oxygen are added

/1,2/. When dissolved sulphur and oxygen are removed from the

liquid, graphite growth will proceed as nodules according to

its crystallographic features shown schematically in Figure

1.

Most commonly, magnesium is applied for nodularizing cast

irons, although cerium, calcium and certain other elements

will also contribute to desulphurizing and deoxidizing.

Normally, magnesium is added to liquid iron either as Mg-

metal or as a magnesium-ferrosilicon alloy (Mg-FeSi). In the

5

following, principal differences between Mg-metal and Mg-FeSi

treatment will be described, and how the respective agents

affect resulting magnesium recovery, reaction product

composition and distribution, and finally inoculation

effectiveness.

3. Principal Nodularizing Reactions

The reactions taking place in liquid iron during magnesium

treatment is described as desulphurizing and deoxidizing

reactions. This means that dissolved sulphur and oxygen are

tied to the nodularizing elements forming some sulphide and

oxide compounds. It is found from several previous studies

that small contents of dissolved S or O may cause a reduction

in graphite nodularity, and consequently these elements must

be effectively removed from dissolution /3,4,5/. This is done

by adding a strong sulphide and oxide former as magnesium.

Magnesium will form sulphides according to the following

reaction:

Mg + S = MgS (s) (1)

MgS is a stable sulphide compound, with an energy of

formation of ?G = -232 kJ and a melting point of 2000°C. This

means that very little sulphur will remain dissolved in the

iron after a proper addition of magnesium. Precipitation and

distribution of such sulphide compounds will depend on type

of magnesium agent applied and the treatment process

involved. For most commercial processes, calcium and cerium

will also be deliberately added in controlled amounts, tying

up some of the sulphur available.

6

Dissolved oxygen present in all commercial base irons must

also be removed to obtain successful spheroidization of the

graphite phase. The amount of dissolved oxygen may vary

significantly according to melting process and charge

materials applied, and consequently the need for nodularizing

elements to tie-up oxygen may also vary. Normally, magnesium

will form oxide or silicate compounds when added to liquid

iron according to the following deoxidizing reactions:

Mg + O = MgO (s) (2) Mg + Si + 3O = MgSiO3 (s) (3)

2Mg + Si + 4O = Mg2SiO4 (s) (4)

Which reaction products, magnesium oxide (MgO), enstatite(MgSiO3) and forsterite (Mg2SiO4), that will dominate depends

on the nodularizing agent applied. Precipitation and

distribution of these reaction products will strongly depend

on type of magnesium agent applied and the treatment process.

Pure Mg-oxide will be the dominant reaction product from Mg-

metal additions, while enstatite and forsterite silicates

dominate from Mg-FeSi alloy additions. Magnesium oxide and

silicates are among the most stable oxide compounds known,

which means that virtually no oxygen will remain dissolved in

the iron after an appropriate magnesium addition. Stability(free energy of formation, ?G) and melting point (Tm) for the

oxide reaction products are:

Magnesium oxide ∆G = -401 kJ Tm = 2830°C

Enstatite ∆G = -1060 kJ Tm = 1580°C

Forsterite ∆G = -1490 kJ Tm = 1900°C

This means that in the presence of appropriate silicon

contents, enstatite and forsterite will be the predominating

7

reaction products. Most Mg-FeSi alloys contain about 45 per

cent silicon, which is found to enhance precipitation of

silicates as enstatite and forsterite.

8

4. Magnesium Treatment Reactivity

Mg-metal are normally added to liquid iron in a converter

process or in a hollow tube (cored wire). Mg-FeSi alloys are

added either in a ladle process, in a continuous reaction

chamber or directly in the mould. A characteristic difference

between a metal addition and a ferroalloy addition is the

reactivity of the respective materials. Inevitably, Mg-metal

addition will give a violent reaction since magnesium has

limited solubility in liquid iron and addition is made at

temperatures above the boiling point of magnesium. Boiling of

magnesium will promote strong turbulence in the liquid and

floating Mg-gas bubbles will effectively remove sulphide and

oxide products (i.e. MgS and MgO) to the top surface.

Accompanied by significant formation of MgO-fume and smoke,

metal addition will result in poor magnesium recovery and

consequently potential hazardous environmental conditions.

Normally, magnesium recoveries from 20 to 50 per cent are

obtained for treatment processes applying magnesium metal

(e.g. converter and cored wire) /6,7,8,9/.

Mg-FeSi alloy additions will, according to the presence of

silicon, calcium and cerium, give less violent reaction in

iron with less effective removal of the smallest reaction

products from the treatment vessel. Silicon in the Mg-FeSi

alloy will increase the magnesium solubility locally during

dissolution, and result in formation of stable Mg-silicates

(enstatite and forsterite) instead of pure MgO. Calcium in

the Mg-FeSi alloy will be present as Ca-Mg intermetallics

which are less reactive compounds than the Mg-Si compounds(e.g. Mg2Si) /10/. Consequently, silicon and calcium reduce

the reaction violence and products will to a large extent

remain as a fine dispersion of particles in the treated iron

9

instead of separating to the top surface. The less violent

treatment will, for the best available processes, give low

fume, smoke and slag formations, and consequently high

magnesium recoveries. Normally, magnesium recoveries from 60

up to 90 per cent are obtainable for the best treatment

processes applying Mg-FeSi alloys (e.g. mould treatment or

tundish cover process) /11,12,13/.

5. Nucleation Potentials from Nodularizing

As described previously, the nodularizing treatment is a

desulphurizing and deoxidizing process. Nevertheless, it is

not necessarily desirable to remove the sulphides and oxides

from the liquid iron to the top slag. To obtain effective

nucleation during the subsequent inoculation it is

advantageous to inhibit the reaction products from

agglomerating and floating, and rather promote a fine

dispersion of micro-particles in the iron. Such micro-

particles will act as potential sites for heterogeneous

graphite nucleation during solidification. Hence, an

effective nodularizing process, which also gives a good basis

for inoculation, is characterized by low delta S and O

values, i.e. the difference between analytical levels of S

and O before and after treatment should preferentially be

low.

A converter or wire treatment process tent to give magnesium

oxide reaction products from the deoxidizing reaction. MgO-

particles have a strong tendency to agglomerate forming

larger slag clusters that float readily, hence giving

effective deoxidation with large quantities of top-slag. A

major disadvantage is that, due to the oxide agglomeration

10

and flotation, a significant part of the potential nuclei

particles will also be removed from the iron. This may cause

problems avoiding carbides in thin section castings even

after adding an effective inoculant material.

Treatment processes applying magnesium-ferrosilicon alloys

give less violent reactions and hence less effective removal

of potential nuclei from the melt. Also, due to interfacial

energy phenomena, magnesium silicates tend to form smaller

particles than magnesium oxide. This means that a larger

fraction of oxides remains as a fine particle dispersion in

the iron after treatment (i.e. magnesium sulphides and

silicates). These particles are very small and highly

numerous. Normal number densities lie in the area of some 100

thousands per cube millimetre with an average size is about

0.5 to 1 micron in diameter /14/. Such small particles will

float very slowly according to Stokes' law, and hence they

will remain in the liquid during handling and pouring. Figure

2 represents schematic segments of liquid iron treated with

Mg-metal and Mg-FeSi, respectively. As will be described

later, particles formed during nodularizing treatment makes

an important basis for the effectiveness of the subsequent

inoculation.

6. Example of an Effective Nodularizing Process

In the present chapter, an example of a very effective

nodularizing process will be described. The process is an

optimized tundish cover ladle treatment applying magnesium-

ferrosilicon alloys. Figure 3 shows a treatment ladle design

for this process. Preferentially, the height:width ratio

should be at least 2:1, with larger ratios improving the

11

process even more. The optimum treatment vessel for such

tundish cover process will be like a long tube divided

vertically for lining maintenance. The tundish cover process

is suitable for all treatment sizes ranging from 100 kgs to

10 tonnes /13/.

The tundish cover lid may be designed to suit a range of

different foundry conditions. Examples are fixed lid,

automatic lift or manual removable lid. The metal outlet

should be sized to suit the amount of iron being treated. If

a divider wall or alloy chamber is applied, as shown in

Figure 3, the metal stream must fall into the part of the

bottom not containing the alloy. By such process, very high

recoveries and good process consistency can easily be

obtained provided the right alloy is applied. Consequently,

low levels of fume and smoke will escape from the vessel,

giving a good foundry environment. A separate smoke

collection system will normally not be required for an

optimized tundish cover ladle.

Magnesium-ferrosilicon alloys suited to the tundish cover

process can be of various composition and sizing. Optimum

alloy choice depends on base iron composition, temperature,

treatment size, etc. Calcium and rare earth contents are

adjusted to reduce reaction violence and neutralize trace

elements in the base iron. Sizing of the alloy is also of

vital importance to its reactivity. Generally, it can be said

that a wide size range will give dense packing in the alloy

chamber (sandwich), resulting in instant agglomeration of

alloy particles when exposed to the heat from the liquid

above. The alloy will then dissolve slowly and gentle

downwards to give a controlled, calm reaction with high

efficiency and good reproducibility. At optimum conditions,

12

magnesium recoveries up to about 85 per cent can easily be

obtained.

Table 1 shows a comparison of some characteristic features

for the tundish cover process compared to the cored wire

process. As can be seen from the Table, both processes show

some positive and some negative characteristics.

7. Residual Magnesium in Ductile Iron

Most ductile iron foundries apply either a spectrograph or

spectrometer to measure the so-called residual magnesium

content. From experience, a certain lower limit of magnesium

is found at each foundry, that provides the required

nodularity in their castings. This limit may vary

significantly from one foundry to another. While one foundry

experience excellent nodularity at 0.025 per cent magnesium,

another may need 0.06 per cent magnesium to obtain good

nodule structures. These differences result from certain

variables that will be explained in the following.

First, the total analytical or residual magnesium content in

ductile iron is comprised of three different contributions.

Both dissolved magnesium, magnesium containing micro-

particles, as well as occasional slag or dross particles will

contribute to the total analytical magnesium found. There

exist no good analytical method separating between these

three contributions, and hence magnesium analyses in

foundries will represent both dissolved Mg as well as Mg-

compounds. The fraction of micro-particles and slags in a

sample will vary significantly from one condition to another,

hence resulting in the variations in analytical magnesium

observed between foundries.

13

Another important factor affecting the analytical magnesium

values is the applied sampling technique. If a sample is

taken from the top of a waiting ladle and from the last

liquid to be poured, differences will be found. Micro-

particles and slag particles will float in the liquid

according to Stokes' law, the larger particles floating much

faster than the smaller. This will result in an inhomogeneous

distribution of non-metallics in the iron during time, and

the analytical magnesium will vary accordingly as a function

of time and sampling position. These phenomena will also

contribute to the variations observed from foundry to

foundry, since time factors and sampling techniques may be

quite different.

Also, flotation of particles explains much of the so-called

fading of magnesium during time. As slag particles float to

the top surface, analytical magnesium in the iron will

inevitably drop. The smaller micro-particles will not

contribute that much to fading, since these float at very

slow velocities. Figure 4 shows flotation velocities of non-

metallics in liquid iron as a function of particle size.

Consequently, fading of magnesium will in fact also be a

positive contribution to the iron cleanliness, since it

represents a removal of slag and dross to the surface. Figure

5 gives a schematic representation of the contributions to

magnesium fading of ductile iron.

The higher magnesium recoveries for some treatment processes

are also to a great extent related to removal of reaction

products. When a tundish cover process is applied,

desulphurizing and deoxidizing products will remain in the

liquid as a fine dispersion of micro-particles hence

14

resulting in high analytical magnesium content in such irons.

This explains the direct relation between a slow and gentle

reaction of magnesium-ferrosilicon, high magnesium recovery,

low sulphur and oxygen removal, and finally a high nucleation

potential.

In fact, inhomogeneous distribution of reaction products and

inconsistent sampling techniques often leads to quite large

variations in analytical magnesium contents from treatment to

treatment for any process applied. Due to this, the real

process consistency may be better than the magnesium analyses

should indicate. This is important to note when the goodness

of a treatment method is being evaluated by means of

reproducibility in magnesium contents.

8. Inoculation Mechanisms of Ductile Iron

In inoculation technology the road between trial and error

approach to a well-founded theoretical one is still

incompletely mapped and obscure. The complexity of the

thermodynamics, kinetics and interfacial phenomena involved

prevents a single, comprehensive theory to be formed from

which a workable alloy recipe can be deducted. Nevertheless,

since cast iron microstructure control is the present key

issue where suitable additions of oxide and sulphide forming

elements to the melt are involved, heterogeneous nucleation

appears as an essential theoretical feature /15/.

Whereas nodularizing, for instance through a magnesium

treatment, is required for graphite spheroidization,

inoculation is a way of controlling microstructure by

minimizing undercooling and increasing the number of graphite

15

nodules during cast iron solidification. Added to the liquid

iron just prior to casting the inoculant provides a suitable

phase for the graphite nodule nucleation upon cooling. The

most prominent inoculants presently used are ferrosilicon

alloys containing small quantities of elements, such as Ca,

Al, Ba, or Sr. The micro-inclusions formed are complex and of

a heterogeneous chemical nature. After nodularizing,

magnesium containing sulphides and silicates can form, and

with reference to Figure 6(a), the dominating constituentphases are MgS, MgO·SiO2 (enstatite) and 2MgO·SiO2(forsterite) /14/.

After inoculation with Ca, Ba or Sr-containing ferrosilicon,hexagonal silicate phases of the XO·Al2O3·2SiO2 or the XO·SiO2type form at the surface of inclusions from nodularization,

Figure 6(b). The presence of phases of this nature will

enhance the nucleating potency of the inclusions with respect

to graphite. It is important to note that inoculation of

ductile iron does not provide formation of new nuclei

particles in the iron, but rather modify the surface of

existing micro-products from the nodularizing treatment. As

illustrated in Figure 7, the (001) basal planes offer

particularly favourable sites for graphite nucleation since

these facets represents a good match for development of

coherent/semi-coherent low energy interfaces between

substrate and nucleus. High purity ferrosilicon doesn't show

an inoculating effect /14/. This highlights the fundamental

importance of the minor elements contained in the alloy, and

in search for more efficient inoculants the recognition of

nucleation theory as a guiding principle should be duly

observed.

16

9. Summary

The following principal effects of nodularizing and

inoculation have been reviewed in this paper:

Nodularizing of ductile iron can be regarded as a

simultaneous desulphurizing and deoxidizing treatment, where

elements having strong affinity to both sulphur and oxygen

are added. Depending on the nodularizing agent applied,

reaction products will be MgS, MgO or complex Mg-silicates.

Magnesium treatment makes an important basis for the

effectiveness of the subsequent inoculation. It is shown that

the majority of micro-inclusions in ductile iron are primary

and secondary products of magnesium treatment. Optimized

nodularizing treatments will give a calm reaction resulting

in a high number of micro-particles which act as potential

nucleation sites for graphite during solidification.

Analytical magnesium content in ductile iron is comprised of

both dissolved magnesium, Mg-containing micro-particles, and

slag or dross particles. Treatment method and sampling

technique are both of vital importance to the analytical

magnesium levels found in commercial ductile irons.

17

10. References

/1/ K.Hurfurth: "Investigations into the influence of

various

additions on the surface tension of liquid cast iron

with

the aim of finding relationships between the surface

tension and the occurrence of various forms of

graphite",

Freiberger Forschungsh., B.105, 1966, p.267.

/2/ R.H.McSwain and C.E.Bates: Surface and interfacial

energy

relationships controlling graphite formation in cast

iron", Proc. 2nd Int. symposium on the metallurgy of

cast

iron, Geneva, 1974, p.423.

/3/ J.P.Sadocha and J.E.Gruzleski: "The mechanism of

graphite

spheroid formation in pure Fe-C-Si alloys", Proc. 2nd

Int.

symposium on the metallurgy of cast iron, Geneva, 1974,

p.443.

/4/ S.C.Clow: "The effect and control of sulfur in cast

iron",

AFS Intl. Cast Metals J., Sept. 1979, p.45.

/5/ R.W.Heine: "Magnesium requirements of ductile iron", AFS

Transactions, 1989, p.485.

18

/6/ H.J.Blicker and D.Michel: "A magnesium wire procedure

for

the manufacture of nodular graphite iron", Castcon'92,

11-

12 June 1992, Stratford-upon-Avon.

/7/ R.Norris, K.Pashley and B.Newton: "Application of

magnesium wire when pouring SG iron from an automatic

pouring furnace", Castcon'92, 11-12 June 1992,

Stratford-

upon-Avon.

/8/ J.Rotella and R.Mickelson: "Using cored wire in the

production of ductile iron", AFS Transactions, 1991,

p.519.

/9/ A.F.Hieber and T.Watmough: "An in-ladle treatment

process

for producing ductile iron with elemental magnesium" AFS

Transactions, 1980, p.289.

/10/ Elkem Magnesium-Ferrosilicon Brochure, April 1989.

/11/ T.L.Forshey, G.E.Isenberg, R.D.Keller,Jr., and

C.R.Loper,Jr.: "Modification of, and production

experience

with, the tundish cover for ductile iron treatment", AFS

Transactions, 1982, p.53.

/12/ R.D.Forrest and H.Wolfensberger: "Improved ladle

treatment

of ductile iron by means of the tundish cover", AFS

Transactions, 1980, p.421.

19

/13/ D.White: "Tundish treatment ladle developments", Arab

foundry symposium, Arabcast'91, Cairo, 7-14 Dec., 1991.

/14/ T.Skaland, Ø.Grong and T.Grong: "A model for the

graphite

formation in ductile cast iron: Part I. Inoculation

mechanisms", Metallurgical Transactions A, Vol.24A,

1993,

p.2321.

/15/ Ø.Grong, T.Grong and T.Skaland: "Principle guidelines

for

new ferroalloy developments", International ferroalloy

conference, Infacon 7, Trondheim, Norway, 11-14 June

1995.

20

List of Captions

Table:

1. Comparison of magnesium treatment processes.

Figures:

1. Graphite growth morphologies.

2. Particle distribution in liquid iron.

3. Tundish cover ladle design.

4. Flotation of inclusion in liquid iron.

5. Fading of magnesium during holding.

6. Non-metallic inclusions in ductile iron.

7. Lattice arrangements after inoculation.

21

Table 1. Comparison of some characteristic features for

magnesium cored wire and tundish cover process.

Cored Wire Tundish Cover

Treatment agent Mg-metal or

MgSi

Mg-FeSi

Magnesium recovery 30 - 50 % 50 - 80 %

Equipment costs Medium Low

Fume emitted Medium to high Low

Restrictions on base

iron sulphur content

No max. 0,03 %

Metal weight

restrictions

> 500 kg No

Inoculation effect

from treatment

Low Medium

Violence of reaction High Low

Possible risk of

excessive Si build-up

No Yes

Size of production

unit suited to

process

Medium to large Small to large

22

Figure 1. Graphite growth morphologies for nodular and flake

graphite /1/.

23

(a) (b)

Figure 2. Schematic representation of particle distribution

in

liquid iron treated with; (a) Mg-metal, (b) Mg-FeSi

alloy.

24

Figure 3. Example of tundish cover ladle design.

25

Figure 4. Flotation of inclusions in liquid iron as a

function

26

of particle size according to Stoke' law /13/.

27

Figure 5. Fading of magnesium during holding of treated

ductile

iron.

28

(a)

(b)

Figure 6. Non-metallic inclusions commonly observed in

ductile

cast iron; (a) Constituent phases present after

nodularization with magnesium ferrosilicon, (b)

Phases formed after inoculation with Ca, Ba or Sr

containing ferrosilicon /14/.

29

(a)

(b)

Figure 7. Details of lattice arrangement at nucleus/substrate

interface after inoculation; (a) Coherentgraphite/BaO·SiO2 interface, (b) Coherent

graphite/CaO·Al2O3·2SiO2 interface /14/.

by George Goodrich

Proper machining of test specimen is crucial for reliable results. GGG 40.3specification requires a "U" notch in the machined bar. Proper machining of the "U" notch specimen isvery critical before testing the impact specimen. The way the "U" notch is machined can affect theresult. The way the "U" notch is machined can affect the result. For example, if the "U" notch isground instead of broaching (even if one uses a new broach tool) the impact values are two to threefoot pounds higher.

In general, the impact value increases as the "U" notch is broached vs. milled vs. ground.

Broaching / Milling / Grinding

--------------------Impact Value Increases

There was no difference found in the impact value of a "U" notch specimen vs. a "V" notchspecimen, under the following conditions:

"V" notch 2mm deep 45o angle 0.25mm radius at the root ofthe notch

"U" notch 3mm deep 1mm radius "U" notch

The fracture appearance of the impact specimen can give some indication of the impace value.Higher amount of crystalline (clevage) white shiny appearance in the fracture will indicate less impact.

Investigation was conducted from samples received from a foundry which experienced widefluctuations in impact values. Above .011% molybdenum, impact was poor.

Conclusions:

1. The type of machining the "U" notch is very critical for reliable accurate impact values.2. Grinding the notch gave the highest impact values.3. Presence of intercellular carbides are detrimental for impact values.4. Carbon, silicon, nickel and molybdenum have the most influential effect on impact value.5. Regression formula of chemistry vs. impact is available in the literature.

The chemistry of the sample had the following range:

Carbon 3.45 - 3.84

Silicon 1.74 - 2.61

Manganese 0.13 - 0.23

Chromium 0.034 - 0.042

Aluminum 0.006 - 0.042

Magnesium 0.041 - 0.080

There was good correlation between silicon and aluminum on the impact values. Higher siliconhad higher aluminum and lower impact.

The foundry was using nickel magnesium alloy for treatment and in-mold inoculation. As inoculation increased, silicon and aluminum nicreasd and impact decreased. Other conclusions on this investigation were:

1. There was good correlation between carbon/silicon/nickel vs. field strength.2. There was good correlation between aluminum/silicon vs. tensile strength.3. There was NO good correlation between aluminum/silicon vs. % elongation.4. There was good correlation of carbon/nickel with impact strength.Note: Analysis of nickel on the sample was not provided.5. When molybdenum was present in the sample, when it was less than .007%, 100% of the barspassed 8 ftl lb. value. 68% of the bars passed 10 ft. lb. value.

Reviewer's Comment (P.H. Mani) Foundries currently producing or planning to produce in future, castings to meet the specification

GGG 40.3 will benefit from this paper on the advantages of grinding the "U" notch to get two or threefoot pounds increased impact values. As opposed to broaching the "U" notch.

The use of looking at fracture appearance is beneficial. But because the impact sample is verysmall, especially at the fractured face, it is difficult to get reliable indications of impact value by lookingat the fracture alone.

Members are advised to refer to DIS Project 19, which correlates fracture appearance with impactvalue on a larger specimen - dynamic tear specimen.

One should use caution in interpreting the correlation between impact properties and the chemistryof the iron. It is the metallurgy of the iron which influences the impact properties, more than thechemistry.

A Review of Common Metallurgical Defects inDuctile Cast Iron

Causes and Cures

C.M.EcobCustomer Services Manager, Elkem AS, Foundry Products Division

AbstractThe objective of this paper is to provide an overview of some of the most commonmetallurgical defects found in the production of ductile cast iron today. The examples shownhave all been determined during the examination of samples in Elkem’s Research facility inNorway.Whilst many foundries recognise the defects, an appreciation of the possible causes, andtherefore cures, is not always apparent. The causes and cures for the different problems areexamined in the paper. Emphasis is made on shrinkage problems, probably the most commonproblem seen by Elkem’s team of service engineers around the world.

IntroductionMetallurgical defects in ductile iron can be very costly to the foundry, not only because thepart has to be remade or rectified, but due to the unfortunate fact that many defects are notrevealed until after the expensive machining stage. Care in the selection of raw materials,good process control in the melting stage and proper metal handling procedures will go a longway to the prevention of defects.Further, a routine for logging and recording of defect occurrences will reveal which are themajor problem areas, allowing for a systematic elimination of the defects.This paper will examine the most common defects, starting with shrinkage. Deterioration ofaffordable steel scrap qualities, use of incorrect inoculants and nodularisers plus the pressuresto get castings out of the door as fast as possible has led to an increase in the incidences ofshrink/porosity related cases seen by Elkem’s team of technical service engineers. Indeed, theductile iron foundry, which truthfully claims not to have shrinkage concerns is the exceptionto the rule.Other common defects may be divided into two basic categories:

-Those related to nodule shape and size, such as compacted graphite structures,exploded and chunky graphite, graphite floatation, spiky graphite and nodule alignment.

-Those related to inclusions/abnormalities within the matrix, such as flake graphitesurfaces, slag inclusions, carbides and gas.These problem areas are described to aid recognition of the defect and causes are discussedtogether with possible cures.

Shrinkage ControlFigure 1 shows a typical sub-surface shrinkage defect. There are many causes of shrinkage inductile iron, experience globally has shown that about 50% of shrinkage defects are related tosand systems, feeding and gating. The other 50% may be attributed to metallurgical factorssuch as carbon equivalent, temperature, inoculation or high magnesium residuals.

Figure 1: Typical sub-surface shrinkage defect with dendrite arms partly covered with graphitesticking out.

When a shrink or porosity is detected in a casting, there are several immediate and simplesteps that can be taken to identify the cause of the problem. Firstly, the geometry of thecasting should be examined to determine whether the location of the defect is close to a sharpradius or a potential hot spot. At the same time, the sand in the region of the shrink should beexamined to look for any soft spots. Sand integrity accounts for a high proportion ofshrinkage defects and a worn seal on the moulding machine, for example, resulting in a lowersand compaction can often be the cause of an unexplained sudden outbreak of shrinkage.

The second avenue of investigation should be the gating / runner designs and the feeding ofthe casting. Whilst many foundries have computer aided design systems, patterns are oftenaltered slightly over the years at shop floor level and can be significantly different from theoriginal design. Also, changes to the feeder specification can lead to different burncharacteristics and metal solidification patterns. This can affect the amounts of feed metalavailable to different parts of the casting.

Metallurgically, there are many factors that can affect the shrinkage tendency. Figure 2 showsthe relationship between magnesium and shrinkage.

Figure 2: Effect of magnesium content on shrinkage

Magnesium, apart from being one of the most powerful carbide stabilisers, has a markedeffect on the shrinkage tendency of ductile irons. Foundries operating at the higher end of themagnesium range, 0.05% or above, will find that the iron is more prone to shrink thanfoundries operating at lower, but very acceptable, levels, say 0.035-0.04%.

Both under-inoculation and over-inoculation can cause shrinkage. In the case of under-inoculation, not enough dissolved carbon is precipitated as graphite. Graphite nodules have afar lower density than the matrix and to precipitate the low density, high volume graphite hasan overall expansion effect, which helps to counter the natural tendency of the iron to shrink.With over-inoculation, too many nucleation points are active early in the solidification,resulting in an early expansion and sometimes large mould wall movements. Later in thesolidification, when feeders become inactive and contraction takes place, there is no graphitecoming out from solution to counteract the contraction and the result is shrinkage between theeutectic cells.

In many foundries, the microstructure shows even sized nodules (accounting for the fact thatthe section cuts through nodules in 2-dimensions). Many foundrymen still consider this to bea good structure, even though the iron is prone to shrinkage. Nodularisers and specialistinoculants are available these days, which help to counter shrinkage by giving a skewednodule distribution. These structures are shown in figure 3.

Figure 3: The same base iron treated with two different nodularisers resulting in a) Skewednodule distribution b) Unskewed nodule distribution

A skewed nodule distribution indicates that some nodules are being created late in thesolidification process and the drawing of graphite from solution at this stage is a veryeffective way to counter shrink. Most inoculants act almost instantaneously and this gives theeven nodule size effect. Once the potency of the inoculant has gone, then there is no driver tocreate nodules late in the solidification and shrinkage can be the result. More recently,nodularisers have been developed by Elkem that have the same effect of producing theskewed and shrink reducing nodule distribution curve.

A low carbon equivalent, or metal that has been held for some time at temperature, due to amechanical breakdown, for example, is also prone to shrinkage. In these cases, the inherentnuclei within the melt will be low and some preconditioning may be necessary to achieve agood level of nucleation.

Compacted Graphite within the structure.Figure 4 shows a good example of compacted graphite in the structure. There are severalcauses of this, the most common being that the nodularisation process has partly failed.Incorrect weighing of the nodulariser or the use of the wrong nodulariser are possible reasonsfor the failure, although a long holding time in the ladle or excessive temperatures can becontributory factors.

Figure 4: Sample with compacted graphite present in the matrix due to partly failednodularisation process.

Another cause of CG particles in the matrix is an incorrect sulphur level in the base iron.Many foundries melt both grey and ductile charges and segregation of returns is essential.During the nodularisation process, the first reactions that take place are a desulphurisation anddeoxidation, these elements combining preferentially with the magnesium. The base sulphurlevel must be accounted for in the calculation of MgFeSi charge weight.A note of caution here with regard to the addition of the MgFeSi to the ladle or treatmentvessel. To add the MgFeSi early to a hot ladle and then hold the ladle for several minutes untilthe moulding line calls for metal is bad practise as the alloy will be burning or oxidising in thebottom of the ladle during this time. Higher and more consistent recoveries can easily beachieved by adding the alloy just before tap from the furnace.

Low Nodule CountAs the compacted graphite mentioned above may commonly be attributed to the nodulariser,then low nodule counts tend to be a function of the inoculant. Figure 5 shows a low countcompared to the foundry’s normal practise. Avoiding long holding times in the furnace andprolonged pouring time post-inoculation will help to achieve consistent nodule counts, as willimproving the responsiveness of the iron via preconditioning. The use of a specialist powerfulinoculant will give the most consistent results.

Figure 5: Two casting with the same metal treatment resulting in a) low nodule count due to longpouring time and b) normal nodule count with normal practise.

Exploded graphiteFigure 6 shows exploded graphite within the structure. Characteristically, exploded graphitelooks exactly as the name might suggest that the graphite has been blown apart.Most MgFeSi alloys contain some rare earth metals, cerium, lanthanum, neodymium,praesodimium etc and these are beneficial in that they neutralise the effects of somedetrimental tramp elements such as lead, bismuth, antimony, titanium etc..Rare earth elementsare also nodularisers and aid the effects of the magnesium. In excess, however, rare earths cancause exploded graphite. This is more especially when high purity charges are used which arelow in tramp elements. Exploded graphite is normally found in thicker section castings withslow cooling rates or at very high carbon equivalent levels.

Figure 6: Sample with exploded graphite present due to excess concentration of rare earthmetals.

Care should be taken when using induction melting as rare earths can be cumulative in theiron. They tend to have very high melting points and do not volatilise, although some will beoxidised and come out in the slag. This is important to note if a low/zero RE containingnodulariser is substitutes to eliminate the problem as it may take time to dilute the residual REout of the system.Should exploded graphite occur, then examination of the rare earth sources should be made –normally the MgFeSi. Melting a virgin charge with steel scrap, pig iron and no returns willquickly show if the returns and/or the MgFeSi are the problem. Latin America and countriesin the Far East tend to use high levels of rare earth in the nodulariser.Reductions in the carbon equivalent may help to reduce exploded graphite.

Chunky graphiteThis is shown in Figure 7. The causes of chunky graphite are exactly the same as for explodedgraphite with the addition that the defect is also found in thinner casting sections and is not assensitive to the carbon equivalent as exploded graphite.

Figure 7: Sample with chunky graphite present due to excess concentration of rare earth metals.

Graphite floatationThis is caused when large, low density graphite nodules are formed during the solidificationof thick section or otherwise slow cooling castings. The nodules, being of a lower densitythan the matrix, tend to float towards the surface of the casting and thus can have a negativeeffect on the mechanical properties (and surface finish) in that region.A reduction in the carbon equivalent will help to control this, as will a reduction in thepouring temperature or increasing the cooling rate of the casting by the use of chills. Theinoculation system should also be examined, as it is likely that the large graphite nodules havebeen formed very early during the solidification process and an inoculant, which will generatemore, smaller nodules, could be an advantage. An example of graphite floatation is shown inFigure 8.

Figure 8: Sample with graphite floatation present due to high carbon equivalent.

Nodule AlignmentFigure 9 shows a classic case of nodule alignment, not too many examples as clear as thishave been seen coming through our laboratory. This is caused by large dendrites growingduring the solidification with the nodules being precipitated between the dendrite arms. Thusthe nodules appear to be aligned. Whilst not normally a serious problem, this can havedetrimental effects on such properties as tensile strength or impact resistance.The normal causes are low carbon equivalent where not enough graphite is precipitatedduring the cooling, under inoculation or too high a pouring temperature.

Figure 9: Sample with nodule alignment caused by large dendrites growing during thesolidification with the nodules being precipitated between the dendrite arms.

Spiky GraphiteThe occurrence of spiky graphite in ductile iron is rare provided that the nodulariser usedcontains a small amount of rare earths. Normally, the rare earth metals neutralise suchelements as lead, bismuth, titanium and antimony, as discussed in the section on explodedgraphite, however the use of a rare earth-free nodulariser where traces of the deleteriouselements are present results in spiky graphite. This is most commonly found in converter ironwhere the separate additions of RE have been left out by human error.The effect of spiky graphite is a dramatic reduction in the mechanical properties of the iron,the spikes provide points of weakness in the structure. Figure 10 shows a typical example ofspiky graphite. The only cure for this type of defect is the addition of rare earths with thenodulariser.

Figure 10 Sample with spiky graphite present in the matrix due to too elevated level of Pb.

Flake Graphite on the Casting Surface

This is commonly seen in foundries, however many ignore the flake graphite on the surface asit forms part of the machining allowance. The defect is illustrated in Figure 11 and clearlyshows the thin layer of flake graphite adjacent to the mould. This is found mainly ingreensand systems and is caused by a build up of sulphur in the sand, which reacts with themagnesium in the iron to form magnesium sulphides and effectively de-nodularise the iron.A higher Mg or Re in the nodulariser can overcome this, subject to shrinkage restrictionsdiscussed earlier, but the most common remedy is to use an inoculant containing cerium. Thishas the effect of re-nodularising the iron locally.

Figure 11 Sample with flake graphite on the surface of the casting due to high sulphur content inthe moulding sand.

CarbidesIn the production of ductile iron, it must be remembered that magnesium is one of the mostpowerful carbide promoters. Coupled with this, the violence of the magnesium reactionduring the nodularisation process tends to destroy nuclei. For these reasons, inoculationrequirements are heavier than for grey irons and under-inoculation or the use of the wronginoculant are amongst the most common causes of chill or carbides in ductile iron.Figure 12 shows typical carbides in a ductile iron structure. Poor inoculation is not the onlycause of carbides, however, and all the potential reasons need to be explored to determine thereason behind carbide formation.

Figure 12 Sample with carbide present in the matrix due to poor inoculation

Steel scrap qualities have already been mentioned in this paper and increasing concentrationsof carbide promoting elements, such as molybdenum, chromium, vanadium etc can lead to thepromotion of carbides. These can be found particularly in the centre of castings or at grainboundaries, where the eutectic solidification front tends to concentrate the elements to thepoint where carbides form. Apart from steel scrap, use of molybdenum containing returns canbe a source of undesirable carbide promoting materials.Low carbon equivalent and high pouring temperatures may also promote carbides,particularly in thin section castings.The cures for carbide problems usually revolve around the use of a more powerful proprietaryinoculant, although nodularisers have been developed which have lower carbide promotingproperties.

Summary

This paper has reviewed the most common metallurgical defects in ductile iron production.Extraneous effects, such as slag and gas have had to be omitted due to space constraints, butthe elimination of these could form a paper on their own.As shrinkage is the most prevalent problem in most ductile foundries, then focus has beenmade on this.Systematic recording of defects, whether found in post casting inspection or even in postfoundry operations is essential to identify the most common and the most costly problemareas. These can then be addressed in order of importance.

NODULAR CAST IRON FATIGUE LIFETIME IN ULTRA-HIGH-CYCLE REGION

František Nový, Otakar Bokůvka, Peter Kopas, Mária Chalupová

University of Zilina in Zilina, Faculty of Mechanical Engineering, Department of Materials Engineering, Univerzitná 1, 010 26 Žilina, Slovak Republic Fax: 00421 41 565 29 40, e-mail: frantisek.novy@fstroj.utc.sk

Introduction The structural materials fatigue damage is dominant one in the threshold state of materials field. Traditionally, the fatigue data, are investigated with testing duration up to N = 107 cycles of loading (steels and cast irons). The fatigue fractures were observed however after billion cycles and more – on the other hand with modern developments in industry, the design lifetime of many components of cars, trains, aircrafts, etc. now often exceeds N = 107 cycles [1-3]. Nodular cast iron and especially Austempered Ductile Iron (ADI) is with regard on the very good mechanical and technological properties very prospective structural material. ADI is used for strongly dynamically loaded machine details e.g. gear and traversing wheels, crankshafts of motor-cars, vans and trucks, rail brakes, pressure pipes in oil industry, in civil and military area [2, 4-6]. Additional, new knowledge about nodular cast iron incl. ADI fatigue lifetime are therefore very important from the point of reliability and safety of used materials. In this paper there are presented original fatigue test results, fatigue lifetime in the ultra-high-cycle region, for nodular cast iron and Austempered Ductile Iron (ADI) which were obtained at high-frequency testing.

Material and testing procedure

The unalloyed cast iron with chemical composition (in weight %) tensile strength Rm and microstructure (pearlite-ferrite), Tab. 1, Fig. 1 was used as a basic experimental material for the study. This nodular cast iron was heat-treated with different austenitization and isothermal transformation procedure in AS 140 salt bath. The aim was to obtain the ADI (Austempered Ductile Iron) with different microstructures respectively tensile strength Rm (Tab. 1). The microstructures of ADI after heat-treatment were characterized by upper or lower bainite, retained austenite and graphite, Fig. 1. For the study were selected experimental physical metallurgy methods-quantitative chemical analysis (chemical composition), metalography analysis (microstructures), tensile test (tensile strength Rm), S.E.M. (fractography) and high-frequency fatigue tests

(fatigue lifetime). High-frequency fatigue tests (HFFT) were carried out with using an ultrasonic testing apparatus KAUP-ZU [2-3, 7-8] for high-frequency sinusoidal cyclic push-pull loading (working frequency f ≈ 20 kHz, temperature T = 20 ± 10 °C, load ratio R = -1, forced specimens cooling with distilled water with anticorrosive inhibitor). Smooth 4-mm-dia round bar specimens polished in the working area by metalography procedures were used (12 to 15 specimens for each testing programme). The fatigue lifetime, stress amplitude vs. number of cycles to failure, was investigated in the region from N ≈ 5.106 cycles to N ≈ 2,5.1010 cycles of loading.

Results and discussion The obtained results, chemical compositions, tensile strength Rm, microstructures, fractography, fatigue lifetime incl. fatigue limit vs. tensile strength Rm and fatigue ratio vs. tensile strength Rm dependence (referred to N = 108, 109 and 1010 cycles of loading) are shown in Tab. 1 resp. in Fig. 1 ÷ Fig. 3. Generally, with regard on the results, see Fig. 1, Fig. 2, Fig. 3 we can say, that fatigue properties of nodular cast iron incl. ADI decreasing in the whole investigated number of cycles loading region. The applied of stress

Tab. 1. Chemical composition, heat-treatment and mechanical properties, nodular cast

iron ( ◊ ) and ADI ( ∆ ● ■ )

C Mn Si P S Cu Ni Mo

3.57 0.97 2.72 0.05 0.022 0.93 0.74 0.037

sign.

heat – treatment austenitization * isothermal

transformation

Rm [MPa]

◊ - 722

∆ 910 °C, 30 min * 380 °C, 60 min 1040

● 920 °C, 30 min * 320 °C, 90 min 1159

■ 920 °C, 30 min * 250 °C, 240 min 1551

22nd DANUBIA-ADRIA Symposium

on Experimental Methods in Solid Mechanics

September 28 - October 1, 2005

MONTICELLI TERME / PARMA - ITALY

amplitude decreasing with the number of cycles increasing, Fig. 1. These specific behaviour of the bainite structures is possible explain with the some factors effect, e.g. content of retained austenite in matrix, plastic properties of matrix, retained austenite transformation to martensite during cycles loading, size of critical defects for fatigue crack initiation, fatigue cracks growth mechanisms, and fatigue crack growth rate and Kath ... [2, 6, 9-11]. These factors have different of intensity effect with regard on the level of transformation temperature and length of isothermal transformation dwell.

The fatigue limit, fatigue ratio (referred to N = 108, 109

and 1010 cycles of loading) decreasing with tensile strength increasing, Fig. 2, Fig. 3. The tensile strength increasing is not accompanied with correspond of fatigue properties increasing. The suitable fatigue properties we can observe in the upper bainite structures with lower tensile strength with compare of lower bainite structures. These facts are with a good agreement with the works carried out at low-frequency fatigue testing and in the region to N = 107 cycles of loading [2, 6, 9]. Conclusions • The fatigue lifetime of nodular cast iron incl. ADI

increases with decreasing stress amplitude continuously in the cycles of number region (5.106 < N < 2,5.1010 cycles).

• The fatigue properties of ADI strongly depends on transformation conditions, e.g. temperature and the length of isothermal transformation dwell.

Acknowledgement This research has been supported by Scientific Grant Agency of Ministry of Education of Slovak Republic and Slovak Academy of Sciences, grant No. 1/1077/04 and is also in the frame of the Joint IT/SK S-T Cooperation Programme Project No. 6NT/SK10 and by the founding from the state program of research and development - New materials and technologies in the construction of machines and equipment, thematic state program “The Development of Personality and Talent of Young Employees and Doctorates of Research and Development under 35 years.” This support is gratefully acknowledged.

References

[1] Mayer H., Stanzl-Tschegg V.: Proc. 2nd Int. Conf. Fatigue in the Very High Cycle Regime, Vienna, A, 2001.

[2] Věchet S., Kohout J., Bokůvka O.: Únavové vlastnosti tvárné litiny, EDIS ŽU Žilina, SK, 2001.

[3] Bokůvka O., Nicoletto G., Kunz L., Palček P., Chalupová M.: Low and High-frequency Fatigue Testing, EDIS ŽU Žilina, SK, 2002.

[4] Dorazil E.: High Strength Austempere Ductile Cast Iron, Academia and Horwood, Praha and Chichester, 1991.

[5] Macko P., Bokůvka O.: Materiálové inžinierstvo, 16, 6, SK, 1999, s. 22-32.

[6] Věchet S.: PhD thesis, VUT-FS-KNoM, Brno, CZ, 1989.

[7] Puškár A.: Vysokofrekvenčná únava materiálov, EDIS ŽU Žilina, SK, 1997.

[8] Puškár A., Bokůvka O., Nicoletto G., Palček P.: Berichte und Informationen, No. 1/97, D, 1997, p. 63.

[9] Věchet S., Švejcar J., Dorazil F.: In. Proc. JSME-MMD Kitakyushu, Japan, p. 249.

[10] Věchet S.: Mechanika 52, No. 217, Opole, PL, 1996, p. 139.

[11] Nový F., Kopas P., Bokůvka O., Chalupová M.: Materiálové inžinierstvo, 3, 10, SK, 2003, s. 191.

0

100

200

300

400

500

600

1,00E+06 1,00E+07 1,00E+08 1,00E+09 1,00E+10 1,00E+11

Number of cycles to failure (N)

Str

ess

amp

litu

de

(MP

a)

As Cast, Rm = 722 MPa

Upper bainite, Rm = 1040 MPa

Lower bainite, Rm = 1159 MPa

Lower bainite, Rm = 1551 MPa

Fig. 1. Fatigue lifetime, nodular cast iron and

ADI, HFFT

0

0,1

0,2

0,3

0,4

0,5

0,6

600 800 1000 1200 1400 1600

Fat

igu

e ra

tio

, σσ σσc

/ Rm

1,00E+08 1,00E+09 1,00E+10

Fig. 2. Fatigue limit vs. tensile strength,

nodular cast iron and ADI, HFFT

0 50

100 150 200 250 300 350 400 450

600 800 1000 1200 1400 1600 Tensile strength, MPa

1,00E+08 1,00E+09 1,00E+10

Fig. 3. Fatigue ratio vs. tensile strength,

nodular cast iron and ADI, HFFT

1

Nodulizing and Inoculation Approachesfor Year 2000 and Beyond - Part 1

by

Dr. R. L. (Rod) NaroASI International, Ltd. – July 30, 2001

Original Paper presented June 15, 2000DIS Meeting, Wichita, Kansas

Abstract: Nodu-Bloc, a new iron-magnesium briquette, offers ductile iron foundries a powerfulalloy that can be used to replace traditional magnesium ferrosilicon (MgFeSi) as well as othermagnesium containing master-alloys. Controlled laboratory tests show that Nodu-Bloc can replaceup to 50 weight percent of MgFeSi. Field trials with Nodu-Bloc confirm these results and showthat Nodu-Bloc replacement of MgFeSi can provide significant cost savings. Foundries convertingto Nodu-Bloc will experience reduced melting costs because less MgFeSi is consumed, less steeland pig iron is required in the charge and far greater levels of foundry returns can be utilized.Foundries can easily save up to $10 or more per ton on molten ductile iron processing costs byincorporating Nodu-Bloc technology.

Introduction: Since the commercialization of ductile iron in 1948, foundries have used numerousmethods to introduce magnesium into molten cast iron. Figure 1 lists some of the approaches andtechniques used over the years. Although some of these processes gained a brief following, andsome have even been used successfully, most have fallen out of favor because of numerousshortcomings. Today, the majority of ductile iron castings made throughout the world are producedusing ladle-metallurgy practices with MgFeSi alloys. It is estimated that MgFeSi alloys are used in65 percent of all ductile irons produced worldwide. In the United States, MgFeSi alloys accountfor an estimated 75 percent of ductile iron production. The remaining ductile iron production ismade using either the magnesium-converter process or magnesium containing wire injection.

During the first decade of the new millennium, ductile iron production is forecast to surpass U.S.

gray iron production, with shipments exceeding 5 million net tons by 2006 (1)

. The supply ofdomestically produced MgFeSi becomes important in assessing whether this important rawmaterial will be available in sufficient quantities to sustain the forecasted growth.

With the International Trade Commission 1999 ruling to rescind dumping duties on ferrosiliconalloys, foreign-produced ferrosilicon alloys have flooded the market, setting near-record low pricesand pushing domestic producers out of the market. See Figure 2. Just last year, American Alloys,a producer of MgFeSi as well as ferrosilicon, was forced into bankruptcy and has closed. Figure 2shows the average production costs of 33 ferrosilicon producers taken from a recent surveyconducted by the Commodities Research Unit, a British economic research firm.

At the present time, all five remaining U.S. ferrosilicon producers are operating at a profit loss.

Their combined, before tax operating income for the last four years is summarized in Table 1.(2)

2

The U.S. producers provided this information at ITC hearings in an unsuccessful attempt to restoredumping duties on ferrosilicon-based products. In addition to increased competition from foreignfirms, the slowing economy and rising energy costs have worsened the plight of domesticproducers. Consequently, some haven't found continued operations to be financially worthwhile.

If, indeed, the supply of domestically manufactured MgFeSi is reduced or curtailed because ofplant closures, alternate nodulizing approaches may be necessary to sustain the projected growth ofductile iron.

Economics of Ferroalloy Production: The U.S. ferroalloy industry was a major market force upuntil the early 1980s. Figure 3 shows U.S. production of ferrosilicon alloys compared with importsfor the time frame 1969 to the present. The decline in production can be linked to several factors.Because of electricity-rate increases, pollution-control costs and strong competition from foreignferrosilicon producers, several domestic producers have gone bankrupt, have closed plants orreduced manufacturing output. Table 2 shows the decline in installed furnace capacity tomanufacture ferrosilicon alloys during the past 20 years.

In the United States, MgFeSi production is dependent on the production of 50 percent ferrosilicon.Fifty percent ferrosilicon is produced in a submerged-arc furnace and then alloyed withmagnesium, calcium and rare earths, also known as mischmetal. The relative cost to produce 50percent ferrosilicon, based on a nominal 25-mil power rate ($0.025/kilowatt), is shown in Figure 4.Electricity and raw materials represent 42.32 percent and 43.92 percent, respectively, of moltenmetal cost at the furnace spout; labor accounts for a modest 8.07 percent of the cost. Major costreductions for producing 50 percent ferrosilicon can only be achieved by renegotiating electricalpower rates. Reducing labor costs has only a minimal effect on overall production costs.

The amount of electricity needed to produce one ton of 50 percent ferrosilicon is 4,500 kilowatts.A single 22-megawatt submerged arc furnace using 25-mil electricity, running 24 hours per day,uses $11,500 of electricity per day, or $4.1 million annually. However, the currentenergy crunch doesn’t bode well for ferrosilicon producers to have access to such low-costelectricity in the future. In fact, during the summer of 2001, at least three plants have curtailedproduction of silicon-based alloys and have sold their contracted electricity back to the powergenerator.

MgFeSi is made by ladle treatment of 50 percent ferrosilicon. Magnesium ingots are plunged intothe ladle, followed by additions of calcium silicon and rare earths. The relative cost to produceMgFeSi is shown in Figure 5. Two ingredients, magnesium ingot and related raw materialsrequired for 50 percent ferrosilicon production account for 70.07 percent of the molten metal costwhile electricity and labor now represent 20.65 percent and 4.14 percent, respectively.

Because electricity has such a significant effect on production costs, foreign ferroalloy producersthat have inexpensive, government-subsidized electricity, have a distinct production-costadvantage. In the survey of ferrosilicon production costs at thirty-three Western World ferrosiliconplants by the Commodities Research Unit, high electricity costs were cited as the reason all U.S.ferrosilicon producers were ranked as high-cost producers. Two-thirds or twenty-four of the

3

ferrosilicon producers surveyed by CRU had lower production costs. All of these overseasproducers had significantly lower power costs.

Currently, only three producers of MgFeSi remain in the United States. Globe Metallurgical,Calvert City Metals and Alloys (CCMA) and Keokuk Ferro-Sil produce MgFeSi as well as othersilicon-based alloys. Keokuk Ferro-Sil, Inc. just started to produce MgFeSi alloys in October 2000

while another, Globe Metallurgical in Beverly, Ohio(3)

, the largest domestic MgFeSi producer, isfor sale. The owner, an investment-holding company, has decided that its return on investment isn'tadequate and that there isn't much hope that market conditions will improve in the near term.Quite simply, there is an excess worldwide capacity to produce silicon-based ferroalloys. Thisoversupply will continue to depress world prices in the foreseeable future. If Globe Metallurgicalis sold, and if the new owner decides to convert the plant to silicon metal production, future U.S.

supplies of MgFeSi will be jeopardized(4)

.

Although considerable production capacity still exists in the United States to manufacture MgFeSi,whether that capacity will be utilized for MgFeSi production remains to be seen. The variousgrades of U.S. ferrosilicon production are shown in Figure 6. It’s apparent that capacity exists toconvert much of the current 50 and 75 percent ferrosilicon production to MgFeSi should the needarise. However, this premise is based on U.S. ferrosilicon producers weathering the continuedonslaught of imports and remaining in business.

U.S. ferrosilicon producers have recently (August 2000) appealed to the International TradeCommission to re-instate dumping duties and restrictions on ferrosilicon imports, but no ruling isexpected soon. Even if a favorable ruling occurs, other non-affected ferrosilicon producingcountries would probably step into the U.S. market. Favorable currency exchange rates and astrong dollar typically are excellent incentives for overseas producers to export ferrosilicon into theU.S. market. Further, there doesn’t appear to be any shortage of ferrosilicon producers who canexport to the U.S.

Without import restrictions, U.S. ferrosilicon production could disappear or be drastically reduced,possibly causing U.S. ductile iron producers to be totally dependent on foreign-produced alloys. Ifthis occurs, the number of available grades and sizes of MgFeSi may be limited. Because oceantransportation is used to ship foreign produced MgFeSi to the United States, it is unlikely thatmultiple grades and sizes would be available because of the logistics problems associated withocean transportation. Only one or two grades of the most commonly used alloy chemistries, of onespecific size, would most likely be available.

MgFeSi Replacement: To meet the growing demand for ductile iron and to circumvent potentialreliance on foreign-produced MgFeSi, progressive foundries need to explore alternate nodulizingmethods. Nodulizing processes that utilize pure magnesium have attracted more attention in recentyears. Eliminating or reducing the amount of silicon based nodulizers has a number of benefits forductile iron producers. Silicon is often an unwanted element and at many foundries, control ofsilicon levels is an economic and technical challenge. High silicon levels typically are the result ofone or more of the following: over-treatment with MgFeSi alloys, improper ladle design, treatmentmethod, treatment temperature and base sulfur level.

4

ASI International, Ltd. has developed a new generation of iron-magnesium alloys (Nodu-Bloc) thataddress potential MgFeSi shortages and as well as provide improved ductile iron silicon control.These new iron-magnesium alloys can reduce or even completely eliminate dependence on MgFeSialloys. The iron-magnesium alloys provide all the cost advantages of pure magnesium processesalong with the ease and forgiving nature of ladle-treatment production techniques. Moreimportantly, by using these low silicon alloys, higher levels of foundry returns can be used in thefurnace charge make-up, resulting in significantly reduced melting costs.

Nodu-Bloc iron-magnesium alloys are manufactured using well-developed powder-metallurgytechniques. Pure magnesium, high-purity iron powder and other additives are carefully blendedand compacted under extremely high pressure. Since a furnace smelting process isn't employed,magnesium levels can consistently be controlled in the range of +/- 0.05 percent. In addition,controlled amounts of calcium, barium, rare earths and copper can easily be incorporated into thebriquettes for those applications requiring special chemistries.

Popular Nodu-Bloc iron-magnesium alloy chemistries are listed below:

Nodu-Bloc Grade 11 - 11% Mg, 0.7% Ca, 0.7% Ba, 3.0% Si, 0.7% C, Balance - IronNodu-Bloc Grade 15 - 15% Mg, 3.0% Ca, 6.0% Si, 2.0% C, Balance - IronNodu-Bloc Grade 20 - 20% Mg, 5.5% Ca, 13.0% Si, 2.0% C, Balance - Iron

Nodu-Bloc briquettes have an almond shape and measure 1.25 inch by 1.0 inch by 1/2 inch, eachhaving a volume of approximately 5 cubic centimeters (see Figure 7). Recently, a somewhat largerpressed disc measuring 4.75 inches in diameter and 1.25 inch thick (350 cc's) and containing either11 percent or 15 percent magnesium has been developed. A schematic of two discs coveringMgFeSi in a ladle bottom is shown in Figure 8. The Nodu-Disc’s have a similar formulation to thesmaller briquettes and can be used as a "reactive cover" material for either iron-magnesium tabletsor standard MgFeSi. The consistent weight of the discs may be advantageous in some applicationswhere weighing charge additions might prove cumbersome. The shape of the pressed disc alsoprovides a more favorable surface area-to-volume ratio, which reduces reactivity in molten iron.

A comparison of Nodu-Bloc iron-magnesium briquettes with 5 percent MgFeSi is shown inTable 3. The density of Nodu-Bloc iron-magnesium briquettes, for a given magnesium level, isconsiderably higher than MgFeSi. However, as with MgFeSi, alloy floatation, especially with the20 percent Nodu-Bloc product, may be a problem. Silicon deficiencies can simply be corrected byadding additional returns to the charge. In many cases, improved foundry return utilization canresult in significantly reduced melting costs.

Although silicon control is necessary in producing high quality ductile iron, many ductile ironfoundries are reluctant to add sufficient returns to their furnace charges for fear of high silicon.Sometimes, these returns are simply sold to scrap dealers at a significant loss. Utilizing thesereturns, in conjunction with Nodu-Bloc replacement of MgFeSi, allows more flexibility in post-inoculation. Higher addition rates of post-inoculants for improved structure and carbide reductioncan now be made while maintaining nominal silicon levels. Higher-base silicon levels fromimproved return utilization will significantly improve refractory live. Lastly, by lowering siliconlevels and precisely controlling these levels, foundries will have improved control over mechanicalproperties such as charpy impacts.

5

Experimental Laboratory Testing and Development: To investigate the effects of various levelsof Nodu-Bloc substitution for MgFeSi, several experimental ductile irons heats were prepared.Three levels of Nodu-Bloc substitution (15%, 30% and 50%) were evaluated as partial replacementfor a nominal 6 percent MgFeSi alloy. The effects of Nodu-Bloc substitution on slag and fumeformation, magnesium recovery, sulfur removal and final microstructure were evaluated during this

laboratory-testing phase (5)

.

Heats of ductile-base iron were prepared in a 2,500 pound induction furnace with the base ironcharge shown below:

220 kilograms (485 pounds) pig iron (Sorel grade)330 kilograms (727.5 pounds) ductile iron returns550 kilograms (1212.5 pounds) steel scrap 25 kilograms recarburizer (55.1 pounds) (crushed electrode grade, 99.9% C, 0.05% S) 8 kilograms (17.63 pounds) 75% Ferrosilicon

Experimental ductile iron treatments were poured into a conventional, 300 kilogram (660 pound)tundish ladle. Two base iron sulfur levels were used, 0.013 percent and 0.033 percent. Treatmenttemperatures were 1,500 oC (2,732 oF), and tundish ladle filling times were 40 seconds. Thetundish ladle had a removable lid and a sandwich divider wall in the ladle bottom. The nominalheight to diameter ratio of the ladle was 2.5-to-1. Nodu-Bloc briquettes containing 21 percentmagnesium were used for the trials along with a 5.9 percent magnesium containing ferrosiliconcontaining 1.0 percent total rare earths. The Nodu-Bloc briquettes were first charged into the ladle.The appropriate amount of MgFeSi was then added as a cover. Finally, 2 kilograms (4.4 pounds)of calcium-bearing 75 percent ferrosilicon was also used as a sandwich cover. Post-inoculationwas accomplished using a 0.30 percent barium-containing ferrosilicon as a stream inoculant in a 68kilogram (150 pound) transfer ladle.

Results for the 0.013% sulfur base iron tests are shown in Table 4. The 15% and 30% Nodu-Blocreplacement levels showed no significant change or reduction in magnesium recovery. However,magnesium recovery for the 50 percent replacement level declined somewhat. The relativelylengthy treatment ladle filling time may have accounted for this reduced recovery.

It was noted during testing that more surface dross was observed at the highest Nodu-Blocreplacement level of 50 percent. It was also noted was that with increasing Nodu-Bloc replacementlevel, treatment reaction intensity increased. Although more flashing and flaring were observed,the overall reaction is best described as being "brighter," not more violent. Since a tundish ladlewas used, the increased reactivity would not be regarded as a problem in a normal tundishoperation. However, with open sandwich ladles, the increased reactivity of 21 percent magnesiumNodu-Bloc could result in some risk of metal splashing. Although not laboratory tested, the 11 and15 percent grades of Nodu-Bloc would provide reduced reactivity.

Test results for the higher 0.033% sulfur base iron are shown in Table 5. For these heats, 0.8 kg(1.76 lbs) of iron pyrites was added to the furnace. The 30% Nodu-Bloc replacement levelsshowed no change in magnesium recovery. However, magnesium recovery for the 50 percentreplacement level declined somewhat. Nodu-Bloc replacement at both the 30% and 50% levels

6

seemed to be much more effective in removing sulfur than 100% MgFeSi additions. Typically,with a 2.5 percent MgFeSi addition to a high (0.033 percent sulfur) base iron, final sulfur levelstypically are above 0.02 percent. In these experiments, the addition of 2.5 percent MgFeSidecreased the base sulfur content from 0.033 percent to 0.023 percent. The 30 percent and 50percent Nodu-Bloc replacement treatments reduced the final sulfur levels to 0.017 percent and0.019 percent. These results tend to indicate that Nodu-Bloc has a somewhat more powerfulcapability to desulfurize a high sulfur base iron compared with just MgFeSi. In more practicalterms, foundries running high base sulfur levels would benefit from using Nodu-Bloc sincenodulizing and desulfurization can both be accomplished without any increase in silicon level.

Microstructural results for the series of experimental treatments are summarized in Table 6. Themicrostructures of all 25 mm section test bars poured with the 0.013% sulfur base iron were all

normal and contained nodule counts ranging from 184 to 237/mm2. Pearlite content was measured

between 60 to 70 percent for all samples. No differences in nodule count or nodularity were notedeven at the highest Nodu-Bloc replacement level. In fact, the 50 percent Nodu-Bloc replacement

showed the highest nodule count (237 N/ mm2 ) and best nodularity even though magnesium

recoveries were somewhat reduced. Similar microstructural results were observed with the 0.033%

sulfur base iron samples; nodule counts ranged from 164 to 178 N/ mm2.

One of the subtler laboratory observations was reduced temperature loss when Nodu-Bloc wasused. For example, at a 30 percent Nodu-Bloc replacement of MgFeSi, the nominal reduction intotal alloy addition rate is 0.30 weight percent. Reducing additions of nodulizing alloys results inless temperature loss from the heating and melting of alloy additions. The heat conservationresulting from 0.30 percent less MgFeSi is estimated to be in the range of 20oC to 30oC (36oF to54oF). Higher levels of Nodu-Bloc replacement would undoubtedly result in additionaltemperature conservation.

Production Results: To date, several foundries have substituted Nodu-Bloc for MgFeSi as anintegral part of their daily production while many others are in the process of evaluating Nodu-Bloc. The production experience of three vastly different ductile iron foundries, each of which haddifferent needs, is discussed in detail in this section.

Foundry A is a medium-sized, high-production foundry producing ductile iron parts for theautomotive and truck industries. Daily production capacity is 280 tons. Although Foundry A has acasting yield which ranges from 45 percent to 55 percent, they generate more returns, in the form ofgates, risers and pouring basins, than they can remelt. They needed an economical way to increasereturns utilization without the accompanying increase in silicon levels. To accomplish these goals,an economical, low-silicon nodulizer needed to be found. Nodu-Bloc 15 met these goals.

Foundry A utilizes three 10-ton induction furnaces for melting. A 2,000 pound capacity open ladlewith a height-to-diameter ratio of 2.5-to-1 is used for ductile iron treatments. Extensive tests withNodu-Bloc iron-magnesium briquettes containing 15 percent magnesium were conducted. It wasfound that a 25 percent Nodu-Bloc replacement, based on total magnesium, allowed the foundry touse an additional 400 pounds of returns per furnace charge and reduce steel scrap levels by anequivalent 400 lbs.

7

Nodulizing is accomplished using the sandwich technique. The appropriate amount of MgFeSi isweighed and placed in a charging container. Next, the Nodu-Bloc iron-magnesium briquettes areplace over the MgFeSi. The charge container is then dumped into a pocket in a completely empty,heated ladle. Foundry grade 75% ferrosilicon is then added to the pocket as additional covermaterial, followed by twelve pounds of cover steel. Residual magnesium levels ranged from 0.035percent to 0.040 percent.

Table 7 shows a comparison of the furnace charge makeup as well as levels of nodulizersemployed prior to and after incorporation of Nodu-Bloc. Little-to-no difference in magnesium flareor reactivity was noted by operating personnel when Node-Bloc was used. The favorable height todiameter dimensions of the sandwich ladle most likely accounted for the modest reaction.

The 25 percent magnesium Nodu-Bloc replacement provided identical microstructural resultscompared to nodulizing with 100 percent MgFeSi. Nodule count, nodularity and matrix structuresremained unchanged. Average nodule count is 275 with an average nodularity rating of 95 percent.Average casting section size is five-eighths of an inch with section sizes ranging between a quarterinch to two inches.

The foundry has realized significant cost savings by utilizing 11.21 percent more returns in thecharge make-up. Production costs have been reduced by $7.45 per net ton. The level of dailysavings achieved by using a combination of Nodu-Bloc, reduced levels of MgFeSi and increasedfoundry returns in the furnace charge is $1,489 daily. Annually, these savings approach $375,000.It should be noted that the level of savings is largely dependent on how the foundry values itsreturns. In this example, the foundry placed a value of $90.00 per ton on its returns. Thus, withthese types of savings, Nodu-Bloc iron-magnesium briquettes have now been incorporated intodaily production. Trials have been run with Nodu-Discs and have produced encouraging results.Additional trials with the discs are scheduled for in the near future.

Foundry B is a much smaller jobbing foundry producing a variety of ductile iron castings. Dailyproduction is about 25 tons. Because of the jobbing nature of their business, optimizing castingyield becomes difficult due to the fluctuating nature of their production schedule. Foundry B meltswith two 4,000-pound induction furnaces.

Twenty percent magnesium containing Nodu-Bloc briquettes were evaluated as a replacement for6% percent MgFeSi for cost-reduction purposes. Foundry B also had a silicon problem and couldnot utilize all of the returns generated. It was often forced to liquidate excess returns by sellingthem to the local scrap yard. This practice had an adverse effect on their balance sheet since itinvolved a significant write-down of assets.

Nodulizing is accomplished in a 750-pound tundish ladle having a height-to-diameter ratio of2-to-1. MgFeSi is first weighed into a charging container. Then Nodu-Bloc 20% iron-magnesiumbriquettes are place over the MgFeSi. The charge container is then dumped into the completelyempty, heated tundish ladle. Foundry-grade 75% ferrosilicon is then placed over the nodulizers.Finally, 22 pounds of cover steel is added to the ladle.

8

Table 8 shows the furnace charge makeup as well as levels of nodulizers employed by Foundry Bboth prior to and after incorporation of Nodu-Bloc. During the foundry trials, no appreciabledifference in magnesium flare or reactivity occurred during the nodulizing operation.

The 46 percent magnesium Nodu-Bloc replacement provided identical microstructural resultscompared with nodulizing with 100 percent MgFeSi. This small foundry has realized significantcost savings by utilizing 10 percent more returns in the charge make-up. Ductile iron productioncosts have been reduced by $10.00 per net ton. The level of daily savings achieved by using acombination of Nodu-Bloc, reduced levels of MgFeSi and increased foundry returns in the furnacecharge is $295 daily. On an annual basis, these savings approached $75,000, which pleasedfoundry management. Needless to say, Nodu-Bloc iron magnesium briquettes have now beenincorporated into daily production.

Foundry C is also a small, jobbing foundry producing mostly ductile iron castings along with grayiron castings. The foundry uses two one-ton induction furnaces for melting. Foundry C’s primeobjective was to reduce ductile iron production costs by eliminating costly nodular grade pig ironand replacing it with its own foundry returns. This foundry, not unlike many other small foundries,tends to over treat their ductile iron with MgFeSi and, consequently, is always battling a siliconproblem. The reasons for over treatment include MgFeSi is used for desulfurization since base ironsulfurs approach 0.02 percent, non-ideal treatment ladle dimensions, and lengthy ladle filling timesdue to the tilting mechanism on the induction furnaces.

Nodulizing is accomplished in a 2,000-pound open ladle using the sandwich process. The height todiameter ratio of the ladle is only 1.25-to-1. The treatment is completely empty and pre-heated.Nodu-Bloc 15% briquettes are added to the ladle first, then MgFeSi is placed over the iron-magnesium briquettes, and finally, one 3-pound Nodu-Disc is added as cover. Lastly, 22 pounds offoundry grade 75% ferrosilicon is placed over the nodulizers for “cover”.

Table 9 shows the furnace charge makeup as well as levels of nodulizers employed by Foundry Cboth prior to and after incorporation of Nodu-Bloc. During the foundry trials, only minordifferences in magnesium flare and reactivity occurred during the nodulizing operation. However,some metal splashing has occurred on an infrequent basis, mainly due to the shallow depth of thetreatment ladle. Residual magnesium levels continued to be in the range of 0.05 to 0.055 percent.

The 57 percent magnesium Nodu-Bloc replacement provided identical microstructural resultscompared with nodulizing with 100 percent MgFeSi. This small foundry has realized significantcost savings by completely eliminating over 1,000 pounds of nodular pig iron from its chargemake-up. Production costs have been reduced by $33.49 per net ton. The level of daily savingsachieved by using a combination of Nodu-Bloc, reduced levels of MgFeSi and increased foundryreturns in the furnace charge is $502 daily. On an annualized basis, these savings are in excess of$126,500. As with Foundries A and B, Nodu-Bloc iron-magnesium briquettes and discs have nowbeen incorporated into daily production.

Discussion: Laboratory testing of Nodu-Bloc replacement for MgFeSi confirmed that it is a viablereplacement for MgFeSi alloys up to 30% substitution. Magnesium recovery and microstructureevaluations showed that Nodu-Bloc replacement was identical to 100% MgFeSi treatment. At

9

higher replacement levels, the nodulizing reaction was more vigorous and some reduction inmagnesium recovery occurred. However, microstructures were identical or slightly better than thelower 30% substitution level. The laboratory findings also suggest that Nodu-Bloc is a more potentdesulfurizer that MgFeSi, particularly when base iron sulfur levels are 0.025 percent and higher.Although the laboratory trials utilized the most potent form of Nodu-Bloc, (21 percent magnesiumcontent), the 11 percent and 15 percent grades would show reduced reactivity.

The summary of production results at three different foundries mostly confirmed the laboratoryfindings. Two of the three foundries used a higher replacement level than 30 percent level andcontinued to produce high-quality ductile iron castings with excellent microstructures. The threecase history foundries all were able to increase their use of ductile iron returns in their charges.The savings levels achieved ranged from $7.50 per ton to over $30.00 per ton. It should be notedthat the savings level calculations greatly depends on what value the foundry places on its ductileiron returns.

Nodu-Bloc replacement of MgFeSi allows foundries to continue to use time-proven ladlemetallurgy practices while also realizing the cost savings of pure magnesium processes. All of thisis achieved without the need for costly wire feeding equipment and alloys or installation of aconverter. Additionally, should supplies of U.S. produced MgFeSi be reduced due to producerplant closings, Nodu-Bloc replacement of MgFeSi is one method to stretch supplies. Additionalresearch work continues to strive for methods that will allow even greater replacement levels ofMgFeSi.

Conclusions:

1.) Extensive laboratory testing of Nodu-Bloc 21% iron-magnesium briquettes has shown that up to30 percent replacement of MgFeSi could be accomplished. Good and comparable magnesiumrecovery and microstructures were obtained from substituting 1.5 weight percent addition rates ofMgFeSi with 1.0 weight percent MgFeSi and 0.10 weight percent Nodu-Bloc. Higher additionrates may result in increased reactivity, possible metal splashing and reduced recoveries, but theseare dependent on ladle design and other foundry variables.

2.) Nodu-Bloc is a very attractive product for silicon control in ductile iron production, since theiron-magnesium briquettes will introduce only trace contributions of silicon to the final castings.This may be of great advantage to foundries producing ferritic ductile iron with requirements forimpact resistance where final silicons of 2.5 percent are often necessary to avoid brittleness.

3.) Production results from three different foundries, showed that Nodu-Bloc replacement ofMgFeSi of up to 50 percent was feasible.

4.) Nodu-Bloc iron-magnesium briquettes appear to be provide greater efficiency in desulfurizationthan MgFeSi in medium sulfur base irons (0.02 to 0.05 percent). In such cases, Nodu-Bloc may bean attractive alternative to competitive treatment processes such as converter and cored wire. Themixture Nodu-Bloc and MgFeSi will still provide the most best advantages of MgFeSi versus puremagnesium when it comes to facilitating good nucleation response of the treated metal.

10

Part II of this paper will address new advances in post-inoculation practices of ductile iron usingnewly developed inoculants that contain a significant amount of oxy-sulfide forming elements.

References:

1.) Modern Castings, January 20002.) Ryan’s Notes, April 12, 19993.) Ryan's Notes, March 19, 20014.) Ryan’s Notes, July 30, 20015.) T. Skland, Elkem Research Laboratory, Norway

Figures 1 through 6 are Powerpoint v. 2000

Figure 1: Ductile Iron Treatment Processes – 50 Years of InnovationFigure 2: U.S. Ferrosilicon Production vs. ImportsFigure 3. 75% Ferrosilicon Average Yearly Price vs. Production CostFigure 4. 50% Ferrosilicon Cost ComponentsFigure 5. Magnesium Ferrosilicon Cost ComponentsFigure 6. U.S. Production of Ferrosilicon AlloysFigure 7: Photograph of Nodu-Bloc iron-magnesium briquettes.Figure 8. Illustration of Nodu-Disc iron-magnesium discs covering magnesium

ferrosilicon.

11

Table 1. “Plight of the U.S. Ferrosilicon Industry”U.S. FeSi Producer Statistics

(International Trade Commission Questionnaire Responses)

2000 est. 1999 est. 1998 actual 1997 actualShipments 180,000 180,000 186,497 189,755(Metric tons of contained silicon for both 50% and 75% Ferrosilicon)

75% FeSi Prices $0.3483* $0.3991 $0.4281 $0.4765(Price per lb. of contained silicon)

Operating Income ($30.7) ($10.6) ($2.8) $15.4(Loss) in Millions

• Average 75% FeSi for Year 2000

Table 2. U.S. FeSi Producer StatisticsFerrosilicon Production - 20 years of contraction

1980 2000 + / - %

Ferrosilicon Producers 7 5No. of Furnaces 39 9 (77%)Installed KVA Capacity 804 224 (72%)Production (Metric Tons) 585,551 408,000 (30.3%)

Table 3. Comparison of Nodu-Bloc Iron-magnesiumBriquettes to MgFeSi

Iron-magnesium MgFeSi Briquettes

Melting Temperature 2,350 to 2,450oF <2,050

oF

Size (typical) 1 in. x 1/4 in. 1.25 x 1.0 x .5 in Magnesium % 3.5% to 11% 11%, 15% and 20%Density 5.5% grade 4.05 grams/cc

11% grade 3.50 grams/cc 4.55 grams/cc 15% grade 4.1 grams/cc 20% grade 3.3 grams/cc

Reactivity in Open Ladle Moderate Moderate & "brighter"Alloy chemistry control Fair Excellentcapability

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Table 4. Research Laboratory Test Results0.013% Sulfur Base-Iron

Mg Substitution % Mg % Sulfur Recovery Tap No. Level#1 - 1.5% Addition 0% - Base 0.042 0.010 48.70% or 9.9 lbs MgFeSi#2 - 8.36 lbs MgFeSi & 15% 0.044 0.009 53% .44 lbs Nodu-Bloc#3 – 7.04 lbs MgFeSi & 30% 0.038 0.007 49% .88 lbs Nodu-Bloc#4 - 4.84 lbs MgFeSi & 50% 0.031% 0.008 40% 1.43 lbs Nodu-Bloc

Notes:1.) Magnesium FeSi alloy – 5.9% Mg, Nodu-Bloc – 21% Mg2.) 300 kg Tundish Ladle, Base Sulfur Level - 0.013%3.) Treatment Temperature - 1,500oC (2,732oF), Tundish filling time - 45 sec.

4.) Post-inoculation - 0.30% Ba containing 75% FeSi stream inoculation into transfer ladle Magnesium Recovery calculations based on the formula:

% Mg recovered = (%) Mg residual + base iron sulfur reduction x 100% (%) % Mg addition

Table 5. Research Laboratory Test Results0.033% Sulfur Base-Iron

Mg Substitution % Mg % Sulfur Recovery Tap No. Level#1 - 2.5% Addition 0% - Base 0.056 0.0235 44%

or 16.5 lbs MgFeSi#2 – 10.67 lbs MgFeSi & 30% 0.055 0.017 47% 1.474 lbs Nodu-Bloc#3 – 8.25 lbs MgFeSi & 50% 0.039 0.019 35% 2.36 lbs Nodu-BlocNotes: 1.) MgFeSi alloy – 5.9% Mg, Nodu-Bloc – 21% Mg

2.) 300 kg Tundish Ladle, Base Sulfur Level - 0.033%3.) Treatment Temperature - 1,500oC (2,732oF), Tundish filling time - 45 sec.

4.) Post-inoculation - 0.30% Ba containing 75% FeSi stream inoculation into transfer ladle 5.) Magnesium Recovery calculations based on the formula:

% Mg recovered = (%) Mg residual + base iron sulfur reduction x 100% (%) % Mg addition

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Table 6. Research Laboratory Microstructure Results0.013% Sulfur Base-Iron – 25 mm Section Size

Nodu-BlocSubstitution Level 0% 15% 30% 50%

Nodule Count (mm2) 184 188 201 237Nodularity % 85% 86% 89% 89%Ferrite Content % 41 42 42 46Pearlite % 59 58 58 54Shape Factor 0.80 0.80 0.81 0.81 Mean Diameter 21.0 21.3 21.2 19.5 (in microns)Notes:1.) Test casting section size - 25 mm

Table 7: Production Experience of Foundry A using 15% Nodu-Bloc Iron-magnesium Briquettes

Original Charge Nodu-Bloc Modified ChargeFoundry Returns 2,100 lbs 2,500 lbsSteel scrap 1,500 lbs 1,100 lbs Carbon 55 lbs 40 lbsSilicon Carbide 4 lbs 4 lbsMgFeSi 27 lbs 21 lbsNodu-Bloc 15% 0 lbs 2.9 lbs 75% Foundry FeSi 11 lbs 11 lbsCover Steel 11 lbs 11 lbs

Final Chemisty% Carbon 3.70% - 3.85% 3.70% - 3.85%% Silicon 2.60% - 2.70% 2.60% - 2.70%% Sulfur 0.007% - 0.009% 0.007% - 0.009%% Magnesium 0.030 - 0.040% 0.03 - 0.040%

Nodule Count (mm2) 275 275

Nodularity 95% 95%Carbides None NoneNotes:1.) 1,900 lb. open ladle, sandwich treatment method

14

Table 8: Production Experience of Foundry B using 20% Nodu-BlocIron-magnesium Briquettes

Original Charge Nodu-Bloc Modified ChargeFoundry Returns 750 lbs 900 lbsSteel scrap 750 lbs 600 lbs Carbon 28 lbs 23 lbsSilicon Carbide 5 lbs 5 lbsMgFeSi 12 lbs 6.5 lbsNodu-Bloc 20% 0 lbs 2.1 lbs Proprietary Inoculant 3.25 lbs75% Foundry FeSi ---- 3.75 lbsCover Steel 22 lbs 22 lbs

Final Chemistry% Carbon 3.60% - 3.75% 3.60% - 3.75%% Silicon 2.50% - 2.65% 2.50% - 2.65%% Sulfur 0.0075% 0.0075%% Magnesium 0.035 - 0.045% 0.035 - 0.045%

Nodule Count (mm2) 225 250

Nodularity 95% 98%Carbides None NoneNotes:1.) 750 lb. tundish treatment ladle

Table 9: Production Experience of Foundry C using 15% Nodu-BlocIron-magnesium Briquettes

Original Charge Nodu-Bloc Modified ChargeFoundry Returns 0 lbs 1,000 lbsSteel scrap 200 lbs 200 lbs Nodular Pig Iron 1,800 lbs 800 lbsCarbon 2 lbs 6 lbs75% FeSi lumps 16 lbs 0 lbsMgFeSi 49 lbs 21 lbsNodu-Bloc 15% 0 lbs 8 lbs Nodu-Disc 15% 0 lbs 3 lbs75% Foundry FeSi 20 lbs 20 lbs

Notes:1.) Base iron sulfur level – 0.025%

15

16

Figure 7. Photograph of Nodu-Bloc iron-magnesium briquettes magnification 1.25x

Figure 8. Illustrations of Nodu-Disc iron magnesium nodulizing discs covering magnesiumferrosilison. Magnification 0.5x

Some Studies of Nodular Graphite Cast Iron*

By Mitsutake ISOBE and Akio CHIDA

The Research Institute for Iron, Steel and Other Metals(Received April 5, 1951)

Synopsis

The experiments were carried out on the formation of nodular graphite in castiron by adding some different kinds of mother alloys to molten cast iron and on theincrease of graphite particles caused by annealing treatment. The yield of modifiedelement in cast iron is increased as the content of alloying element in mother alloyor the amount of mother alloy added to molten cast iron decreases. By using many-components alloy the structure in which the ground mass consisted fully of ferrite,is obtained. The hardness of modified cast iron decreases rapidly by annealing orwith the increase of the diameter of specimen.

I. IntroductionNodularizing of graphite in cast iron by adding cerium was successfully performedby Morrogh and other iñvestigators(’)(2)(3).4)• at the Association of Cast Iron Foundry-men in England during the Second World War. After that, in the U. S. A.(5) theinvestigation was done for the purpose of making the nodular graphite cast iron bymeans of the niagnesium treatment.Here we studied on the conditions of optimum treatment of magnesium added tomother alloys for noclularizing the raphite in cast iron of a certain compositionand confirmed the yield of that mother alloys in a casting practice and also theeffect of annealing upon the casting structure.

II. Experimental method

Throughout the experiment, 35 K. V. A. high frequency induction furnace was usedfor melting the alloys and the temperature of melt Was ‘measured by means of theoptical-pyrometer which is previously well revised. Maximum heating temperatureof melt, casting temperature, and temperature at which the mother alloy was addedinto the melt Were 1400, 1300 and 1340°C respectively. Throughout this study, allthe conditions of experiment was kept always constant. The mother alloy series,Al—Mg, FeESi-Ni--Mg and Fe-Si-Ni-Al-Mg, were made in additic*i to the generallyused Cu-Mg and Ni-Mg alloys, and the yield of respective mother alloys and an

* The 633 rd report of the Research Institute for Iron, Steel and Other Metals.(1) H. Morrogh, W. J. Williams; Jour. Iron & Steel Inst., 155 (1947), 321.(2) H. Morrogh, W. J. Williams; Jour. Iron & Steel Inst., 158 (1948), 306.(3) H. Morrogh, J. W. Grant; Proc. Inst. Brit. Foundrymen, 41 (1947—48, A—29.(4) H. Morrogh; Iron Age, 163 (1948), May 20, 82.(5) International Ni Co., Foundry Trade Journal, 84 (1948), 463.

298 Mitsutake ISOBE and Akio CHIDA

influence of annealing for castings caused by the change of amount of magnesiumwere determined. Both green and dry sand moulds were used.

III. Material used

Table 1 shows the chemical composition of cast iron’ used in this experimentThese white pig irons contain high carbon, low sulphur and low phosphorous.Table 2’ shows the classification and respective chemical compositions of the addedmother alloys.

Table 1. Chemical composition of cast iron in percent.Base Material (1)

C P ‘ S Mn Si

4.01 0.04 0.02 0.058 0.077

Base Material (2)

C P S Mn Si4.21 0.058 0.028 0.025 0.073

Table 2. Classification and respective chemicalcomposition of mother alloys.

Mark Mother alloy Relative content of respective element.A Cu—Mg 75:25

B Ni—Mg ‘ 75:25C D 70:30D F) 80:20

E Al—Mg 90:10

F Fe—Si—Ni—Mg 20 : 50: 10 : 20

G Fe—Si—Ni—Al—Mg 15 : 50: 10: 5 : 20

IV. Results

Table 3 shows some examples of this experiment As shown in Table 3, in thecase of an adding amount of mother alloy is little the yield of magnesium contentof mother alloys to the castings are good, and as the content of magnesium in mo­ther alloys is low, it will be seen the increased ratio of magnesium content in thecastings.Generally speaking, the yield of magnesium extends from 5 to 16 percent, themean value being 8 to 9 percent.Fig. 1 shows the relation between the yield of magnesium in castings and themother alloys added with magnesium which consist of 80:20, 75:25, 70: 30, in thecase of Ni-Mg alloys and 75:25 in the case of Cu-Mg alloy’ respectively and it willbe seen that there are some relationships amOng them.

Table 3. Some examples of this experiment.

Rum Mother Sand Mg % T.C Si Mn S PInoculation

No. alloy mould Added Analysis Yield % % % %I- 1 Cu—Mg green 0.4 % 0.25 0.029 11.6 3.78 2.27 . 0.21 0.027 0.051

75:25 Si— 2 ii 0.50 0.048 9.6 3.80 2.38 0.23 0M26 0.053— 3 p Li 9 0.075 0.011 14.6 3.82 2.41 0.28 0.028 0.048— 6 0.75 0.063 8.4 3.72 - 2.50 0.22 0.027 0.051— 7 ii dry

*0.75 0.070 9.33 3.71 2.36 0.26 0.022 0.052

— 8 LI green II 1.25 0.073 5.84 3.82 2.40 0.29 0.026 0.053— 9 p ii Li 1.00 0.072 7.2 3.89 2.34 0.27 0.029 0.052—10 ‘1 ii 1.50 — — 3.81 2.30 — 0.023 0.051

II— 1 Ni—Mg green 0.4 % 0.25 0.044 17.6 3.76 2.40 0.23 0.027 - 0.05075:25 Si

— 2 ii p p 0.50 0.067 13.4 3.93 2.48 0.20 0.030 0.049— 3 Li Li ii 0.75 0.051 6.8 3.82 2.39 0.27 0.027 0.051— 4 p p p 1.00 0.053 5.3 4.00 2.45 0.25 0.026 0.053— 5 ii p 1.25 0.074 5.92 3.86 2.12 0.31 0.031 0.062— 6 dry -Li 1.50 0.075 5.0 3.64 1.96 0.21 0.023 0.048—10 p - p 0.75 0.073 9.74 3.90 2.33 0.24 0.022 0.050—11 p p Li 0.50 0.069 13.8 3.87 2.34 0.22 0.023 0.047—12 Li ii 1.00 0.057 5.7 3.88 - 2.41 0.23 0.026 0.043

III— 1 - (F) green — 0.80 0.079 9.88 3.59 - 2.05 0.30 0.022 0.039alloy

— 2 P — 0.60 0.070 11.7 3.77 1.81 0.33 0.020 0.039— 3 0.4 % 0.60 0.073 12.1 3.97 2.31 0.26 0.025 0.046

Si— 4 (G) dry — 0.80 0.080 10.0 4.02 2.10 0.24 0.024 0.053

alloy— ii green — 0.80 0.076 9.5 3.99 2.03 0.28 0.028 0.056

IV— 1 Ni—Mg green 0.4 % 0.60 0.082 13.7 3.97 2.41 0.22 0.029 0.05780:20 Si

2 Li Li 9 0.80 0.092 11.5 3.93 2.39 0.21 0.026 0.050— 3 II Li if 0.40 0.079 19.74 3.90 2.37 0.25 0.027 0.051— 4 Li Li P 1.20 0.098 8.17 3.96 2.40 0.27 0.023 0.049V— 1 Ni—Mg green II 0.30 0.042 14.0 3.79 2.37 0.22 0.026 0.052

70:30— 2 Li P 9 0.60 0.062 10.0 3.82 2.33 0.21 0.025 - 0.054— 3 P Li 0.90 0.060 6.66 3.78 2.34 0.23 0.022 0.049— 4 1) Li ii - 1.20 0.063 5.25 3.88 2.36 0.26 0.023 0.050

VI—lO Al—Mg green 9 0.4 0.062 15.5 3.99 2.40 0.20 0.029 0.05390:10

- - —

-

—0—Ni— Mg( 70:30),, (755)“(80:20)

“Cu—Mg(75:25)

0.2 0.3 O..4 Q50.60.7 0.8 0.9 1.0 1.1 12141.4 1.Mg °/ added

Fig. 1. Relation between the amount of added magnesium and its yield.

Myskowsky— Au?horsOflakygraphjje‘Fe3C± Nodniar graphite• Nodniar graphitef1aky+ Noclntar graphite

300 Mitsutake ISOBE and Akio CHIDA

21.020.0

19.018.0

170

16.015.014.0

13.0

120 \11 0

908.0-7Q

6.05.0

/0.14013

01:2011

0.69-

xl

10.08 11

0.07 s •‘ •

0.05 x_.- :004 .

0.03 .7 8

—-8 8

002 -

0 0001 0 0

°0

°O.6 0.8 1.0 1.21.4 1.6 1.8 2,0 2.2 2.4 2:6 2.8 3.0Si%

Fig. 2. Mode of occurrence of graphite.

Some Studies of Nodular Graphite cast Iron 301

b the case of norlular graphite cast iron, Tanimura(6) and Myskowsky(7 studiedon the relationship between the yield of magnesium content and the amount ofsilicon content in the cast iron containing carbon from 2 to 3.7 per cent and re­cogni-ed that the amount of remained magnesium in cast iron is 0.06 percent. But,the amount of carbon con-tent in the cast iron used 280in this experiment extends ‘ofrom 3.8 to 4.2 percent, and 260thus the percentage of car- 250bon in cast iron is some-

-

240what greater than the pre- 230 •

vious investigators.. 220

As shown in Fig. 2, the 210occurrence of nodular 200graphite particles in castiron can be well recogniz­ed from the amount of 0.05 -.

170percent of remamed magne.sium. That is, the occur- 0.2 0.50 0.75 1.00

Mg%ad4edrence of nodular graphite Fig. 3. Relation between the hardness and the amount ofparticles was confirmed added magnesium (green sand mould).even the amount of remain­ed magnesium in castingsis small, when the amountof carbon in cast iron is 280

high compared with the 270

results of previous experi- 260

ments. t250

-Fig. 3 shows the relation 40 7’between the amount of c. 230 C.)

magnesium and the hard. 220

ness of castings by the 210green sand mould. As 2O0 ‘‘

clearly shown in the Fig., 190 .

the hardness of castings - ‘mo •

made from the Cu-Mg ‘170(75: 25) mother alloy are 160higher ‘than that from the 0 10 20, 50 60 70 80 90 100> iUr’Ni-Mg mother alloy. The Mg%, analysisBrinell hardness number Fig. 4. Relation between the hardness and the amount ofof castings treated with magnesium in cast iron.

(6) H. Tanimura; The Iron & Steel Institute of Japan, (1949) Autumn Meeting.(7) E. T. Myskowsky, etc.; Iron Age, 164 (1949), Sept. 8, 78.

302 Mitsutake ISOBE and Akio CHIDA

Cu-Mg mother alloy series indicates the value from 202 to 280, while in the case ofNi—Mg mother alloy it shows the value from 175 to 250. In both cases,, it will beseen that the increased hardness of cast iron accompanied with the increased amountof added magnesium.Fig. 4 shows the relation between the amount ofhardness of castings.

260250240230220 -

210• 200j190q180

x17016010

iron gradually decreases as th diameter of cast specimenextent of variation of diameter.Next, the microscopic investigation will be shown. At a comparatively low treat­ing temperature, the tendency of appearance of ledeburite structure was conspicuous,,but in the case of a green sand mould casting it amount was comparatively small.Photos. 1 and 2 show the micro-structures of. 0.759 Mg cast iron cast in greenand dry sand moulds respectively. In both cases the Ni-Mg (75 : 25) alloy were usedfor the adding mother alloy.In the case of the dry sand mould casting the so-called “bull’s eye” structure i. e.,the graphite particis surrounded by ferrite were appeared, the ground mass beingpearlite. This corresponds to the high ductility iron in C. K. Donoho’s’°) classifica­tion. In the case of the green sand mould castings, the “bull’s eye” structure arecomparatively small and are studded with small particles of cementite in the groundmass of pearlite. Photos. 3, 4, 5 and 6 show the variations of the structure ofmagnesium bearing cast iron which was treated with Ni-Mg mother alloy. Theamount of magnesium were 0.5, LO, 1.25 and 1.5 percent respectively. As the amount

(8) Product Engineering, 170 (1950) 520; E. K. Smith, Iron and Ste’el, 23 (1950) 46,M. G. Fontana,; Industrial and Engineering Chemistry, 42 (1950), 79.

(9) T. Sato; Kinzoku, (Metals), 20 (1950), 13.(10) C. K. Donoho; Iron Age, 164 (1949), Feb. 24, 97.

‘C

—

———-Ni—Mg(75 : 25)Cu—Mg(75:25)

1.0%Mg

O.75%MgO.5%Mg’1’O%Mg-O.25%Mg0.75 %Mg

0.5% M8’xO.25%Mg

remained magnesium and the

Fig. 5 shows the relationbetween the hardness andthe thickness of castingswith various amounts ofmagnesium where the fulllines denote the results ofhardnçss measurement ofcast’ iron to which Ni-Mgalloys were added as amother alloys and thedotted lines are for Cu-Mgalloys. Several data(8.(9concerning the mass effectof the nodular cast ironwere reported, but as wil{be seen from Fig. 5, thehardness of a nodular castincreases, within the small

15’ 20, 25Diameter of cast specimen(mm)

Fig. 5. Relation between the hardness and the masseffect of nodular cast iron.

Some Studies of Nodular Graphite Cast Iron 303

of magnesium increases, the structure containing a hard constitution giadually ap­pears. Photo. 6 shows the ledeburite structure of iS percent magnesium treated castiron. All the specimens mentioned above were cast in the green sand. mould, thediameter of castings being 15mm.A comparatively good structure of nodular cast iron was obtained by the additionof magnesium ranging from 0.5 to 0.8 percent.Table 4 shows the number of graphite particles existing in the different part ofthe ‘specimen, in which A, B, C and D denote the averaged number of graphiteparticles which were observed in each area of 0.53 mm2, at the central portion, 3 mm,5 mm distant from the center, and the portion of the surface of specimen respectively.

Table 4. Number of nodular graphite particles in as-cast specimen(number in area of 0.53 mm2)

Run No. 1—9 11—2 11—3 11—4 11—5 11—6 111—1 1112

Mg % added 1.00 0.50 0.75 1.00 1.25 1.50 0.80 0.60

Number of nodulargraphite particles, 42 30 52 43 50 33 54 79

portion A

if , z’ B 42 27 54 45 47 35 57 85

1 ,

‘C 40 22 59 47 49 36 60 98

if , Fl D 55 37 62 52 59 42 65 112

reioerjn 44 29 56 46 51 36 59 93

From the results of these microscopic observations, the number of nodular graphiteparticles were not very different from portion to portion of the specimen, except thatthe number of particles in the edges was more or less in excess generally. Moreover,particles of nodular graphite existed more in the cast iron treated with a co-ternaryor many-component mother alloy than in that with a binary alloy.

8Oif— 3 0.75 %Mg added

70.if— 5 1.25 % Mg added

60 - -- -- ff1— 1 0.8% Mg added

°. T if— 4 1.00%Mg adder.. I.j5Oj

15 0 45 60Annealing time (mm.)

Fig. 6. The effect of annealing time upon the number ofnodular graphite particles.

I

304 . Mitsutake ISOBE and Akio CHIDA

Table 5 shows the effect of annealing upon the number of nodular graphite particlesin the specimen annealed at 750°C for 1 hour. Generally, the number of nodulargraphite particles was increased distinctly by annealing.

- Table 5. The effect of annealing upon the nodular graphite particlesin the specimen annealed at 750°C for 1 hr.

Run No. 11—3 11—4 11—5 Ill—i

Mg% added 0.75 1.00 1.25 0.80

Number of nodular graphite 52 43 50 54particles (as-cast)

Ap (annealed) 80 53 58 62

Number of particles increa6ed 28 10 8 8Degree of particles increased % 53.8 23.2 16.0 14.8

Number of nodular graphite 54 45 47 57particles (as-cast)

B‘

(annealed) 79 59 60 60Number of particles increased 25 14 13 3Degree of particles increased % 46.2 31.3 27.6 5.26

Number of nodular graphite 59 47 49 60particles Las-cast)

c p (annealed) 76 57 63 63Number of particles increased 17 10 14 3Degree of particles increased % 28.8 21.2 28,5 5.00

Number of nodular graphite 62 52 59 65particles (as.cast)

D (annealed) . 83 64 72 64Number of particles increased 21 12 13 —1Degree of particles increased % 33.8 23.0 22.0 —1.5

Number of nodular graphite 56 46 51 59Mean particles (as-cast)

(annealed) 80 58 64 62value Number of particles increased 24 11 12 3

Degree of particles increased % 41.2 23.9 23.5 5.00

Fig. 6 shows the effect of annealing time upon the number of nodular graphiteparticles. The number increases as the annealing time increases except for the co­ternary mother alloy (F).Fig. 7 shows the relation between the hardness and the annealing time of sandmould castings treated with 0.75 % magnesium of Ni-Mg mother alloy and annealedat 750° for 15, 30, 60 and 80 min. respectively. The hardness rapidly decreases asthe time increases until the variation of hardness is gone at annealing time of60mm.Photos. 7. 8 and 9 show the structure of the same specimen as in photo. 1, annealedat 750’C for 15, 30 and 60 mm. respectively, the ferrite area increases followed bythe extinction of the pearlite area. • In the specimen annealed at 750°C for 15 mm.

Some Studies of Nodular Graphite Cast Jro,z 305

0.7 5% Mg cast iron cast in dry sandmould, treated with Ni-Mg (75 25) motheralloy, etched with picral.

306 Mitsutake 1SOBE and tkio CHIDA

Some Studies of Nodular Graphite Cast Iron 307

the decompositions of the pearliteare in transit; in the one annealed 98 228for 30 mm., the most part of pearlite 2area shows decomposition; in the one

-‘

annealed for 60 mm., the complete•ferrite structure caused by a completedecomposition of pearlite appears.

.. 1

-

From the results above mentioned86

it is known that the velocity of de­composition of pearlite remained inthe nodular graphite cast iron take

‘

place comparatively rapidly.78 146

Photos. 10 and 11 show the as caststructure treated with co-ternary 7mother alloy (F), where the groundmass are wholly ferrite like the an- . 131

nealed specimen just above explained. 0 15 30 - 60 75 90Armealzng timç( mrn..’

Fig. 7 The relation between the hardnessand the annealing times at 750°C

Summary

1. A’ study of graphitization was made by using the white pig iron containinghigh carbon, low phosphorous and, low sulphur.2. As the adding mother alloys, many-cc*ilponents alloys were used simultaneouslywith the commonly used binary alloys.3. By using these mother alloys the cast structure with fully ferrite ground masswas obtained.4. The yield of modified element to cast iron increases as the content of alloyingelement in mother alloy or the amount .of addition of mother alloy to molten irondecreases.

5. By annealing, the hardness of cast iron decreases rapidly.6. The decomposition velocity of pearlite in nodular graphite cast, iron is comp­aratively rapid.7. The hardness of castings decreases as the diametec of specimen increases.8. Structural diagram concerning the shape of graphite and the amount 5of re­tained magnesium and silicon in cast iron was determined and clarified the relation­ships among them, modifing the Myskowsky’s diagram.

Thermal Analysis

Suggestions for Improved Reliability in Thermal Analysis of Cast Irons

1. Liquid Iron Sampling - uniformity & consistency are mandatory

a. Make sure the sample ladle is free of iron and slag and is as hot as practical.

b. Don't use graphite or clay-graphite ladles for sampling since they tend to produce some inoculating effect.

c. Make sure the iron being sampled is as close to the same temperature as possible from sample to sample. (Note that this temperature will vary depending on whether it is a furnace or ladle sample.)

d. For electric furnace samples, best results are usually obtained once a temperature of ~2700

oF(1482

oC) has been reached in the

furnace. e. If sample ladle is filled from another ladle, we recommend filling the

sample ladle, emptying it and refilling before pouring your sample. f. Make sure the cup stand is kept as clean as possible, free from

resin build-up and spilled iron, out of any cold drafts and at as constant a temperature as possible.

g. Make sure the cup and stand are level, not tipped, and that the cup is poured completely full.

h. If frequent samples are poured (more than ~4-5/hr.), use a second stand and alternate between stands to prevent overheating of the stand.

i. Good sampling practice will result in consistent maximum temperature at least 50

oF(10

oC) to 100

oF(38

oC) above the liquidus

temperature. j. When pouring final iron samples, allow at least two minutes after

inoculant addition before pouring your sample. k. Remove the sample from the cup stand as soon as analysis is

completed to minimize resin build-up and to permit the stand to cool prior to the next test.

l. Regularly clean the cup stand and contacts with a wire brush or other means.

2. Equipment Calibration

• from the cup all the way back to the instrument - extremely important!! (Be aware of variations between IPTS thermocouple standards. Cups for the N. American foundry industry are 1948 IPTS. Do you know whether your instrument is being calibrated per 1948, 1968, or 1990 IPTS?)

Table showing errors between IPTS 68 and IPTS 48 and also the temperature differences which exist if IPTS 48 thermocouples are used with IPTS 68 instruments.

(From 9/23/81 memo from L.R. Jones/E-N to W.F. Shaw)

Temp oF (

oC) Instrument 68/TC 48

IPTS 68 IPTS 48 Temp oF T

oF

2000 (1093) 1997.3 (1092) 1995.5 (1091) 4.5

2100 (1149) 2097.1 (1147) 2095.4 (1146) 4.6

2200 (1204) 2196.9 (1203) 2195.1 (1202) 4.9

2300 (1260) 2296.8 (1258) 2294.9 (1257) 5.1

2400 (1316) 2396.7 (1314) 2394.0 (1312) 6.0

2500 (1371) 2496.5 (1369) 2492.5 (1367) 7.5

3. Cup purchasing, storage & monitoring

a. Purchase cups in large quantities (obviously based on your usage) with the specification that all cups in each order be from the same manufacturing lot.

b. Store cups in as warn and dry an atmosphere as possible. If stored in a cold warehouse, make sure cups are brought up to room temperature well in advance of their usage.

c. Before using a new lot number, compare TA curves from current vs. new lot of cups to ensure that no significant differences occur.

d. Include the supplier's lot number on your melt records & note when a change in lot number occurs.

54 FTJ March 2004

Transport

Austempered ductile iron (ADI) is stronger perunit weight than aluminium, highly wearresistant and easier to machine than freemachining steel - and it has the potential forup to 50% cost savings.

Austempering is a special isothermal heattreatment that can be applied to ferrousmaterials for increased strength and toughness.Fig. 1 shows a schematic isothermal diagramcomparing the austempering (green line) andthe quench and tempering (red line) processes.Austempering consists of austenitisingfollowed by rapidly quenching to a temperaturerange (260-385˚C) where the materialtransforms isothermally to form eitherausferrite in cast iron or bainite in steel.

The quench and temper process consistsof austenitising and then rapidly quenchingbelow the martensite start line. The martensiteformed is a very hard, brittle phase and requiresone or more tempering processes to obtainstrength and toughness.

As an isothermal process, austemperingoffers quality advantages versus quench andtemper. Formation of ausferrite and bainiteoccurs over minutes or hours at a singletemperature; dimensional tolerances are morereadily maintained and cracking avoided. Incontrast the formation of martensite occurswith shear immediately as the metaltemperature reaches the martensite start line.Since cooling rates vary according to section,the transformation is not homogeneous,significantly increasing the risks of distortionand cracking.

Austempering services are now availablefrom contractors such as ADI Treatments Ltd(West Bromwich, UK). Working with theEuropean foundry industry and castings users,the company operates large scale, specialistfurnaces and assists in design and developmentof ADI components.

ADI grades, properties and benefitsADI materials are versatile; table 1 listing thegrades in common use. The designer can selectcast composition and heat treatments toprovide specific properties required for theapplication. Fig. 2 illustrates the ausferritemicrostructure, a mix of acicular ferrite andcarbon stabilised austenite that gives ADI itsunique properties.

Cost advantagesThe price of ADI material is lower per kilo thansteel or aluminium, but this accounts for onlya fraction of the potential savings as an ADIdesigned component can save cost at eachstage of manufacture. ADI equivalents can thenbe produced for less than a steel forging or athalf the cost of aluminium parts. Severalfactors favour ADI in value engineering:

• Excellent castability: Readily cast intocomplex shapes, ductile iron has a very highyield rate i.e. the proportion of metalpoured versus metal shipped.

• Lower machining cost: Well suited to nearnet shape casting, ADI requires less startingmaterial and less metal removal. Prior toaustempering, ductile iron exhibits bettermachinability than free machining steels.

Both ductile iron and ADI produce dense,discontinuous chips that are easily handled,further reducing cost per kilo.

• Heat treatment savings: Austemperinggenerally costs less than carburising orinduction hardening, and produces a higherdegree of uniformity and predictabledimensional changes.

• Low energy content: Producing a typicalADI casting consumes 50% less energythan a steel casting, and 80% less energythan a steel forging.

• Best value: When comparing relative costper unit of yield strength, ADI is generallythe best buy.

Performance advantagesDue to superior performance ADI castings arerapidly displacing steel forgings, weldedfabrications, carburised steel, and aluminium inkey applications:

• Strength comparable to steel: Because of itsequivalent strength, nearly 80% of all castand forged steels can be replaced with somegrade of ductile iron or ADI.

• Lower density than steel: The relativeweight per unit of strength of ADI allowseconomy in design without loss ofperformance. For a given shape, an ADIcomponent will be 10% lighter than steel.

• ‘Lighter’ than aluminium: ADI is threetimes stronger than the best cast or forgedaluminium and weighs only 2.5 times asmuch. Because it is twice as stiff, a properlydesigned ADI part can replace aluminium ata weight saving.

• Excellent fatigue strength: ADI’s dynamic

When ductile iron is subjected to anaustempering process, the material undergoesa remarkable transformation - ideal for manyautomotive applications.

ADI solutions aid vehicle design

Fig. 6 Ford Mustang Cobra suspension arm wins casting award (courtesy Intermet and BentelerCorporations)

Grade Min tensile stress Min yield stress Elongation HardnessMPa MPa %

EN-GJS-800-8 800 500 8 260/320

EN-GJS-1000-5 1000 700 5 300/360

EN-GJS-1200-2 1200 850 2 340/440

EN-GJS-1400-1 1400 1100 1 380/480

Table 1. EN1564: 1997 ADI European grades

Fig. 1 Schematic isothermal diagram illustrating theaustempering (green line) plus quench and tempering (redline) processes

FTJ March 2004 55

Transport

materials, TVR created curves in the body thatcould not be accomplished with regular steelstampings. Less visible, the ADI crankshaft isanother significant innovation.

Forged steel was originally chosen as thecrankshaft material for use in the inline sixcylinder engine. Due to high cost ofmanufacture, steel was later deselected andTVR turned to 800/2 ductile iron for the designprototype. Tests were carried out on a benchdynamometer and in vehicles; however theparts failed, in some cases with a fatigue crackat a fillet radius on the flywheel end.

ADI (fig. 4) became the next choice. TVRwas already using the material in V8 enginesbut had initial reservations about distortion inthe manufacture of the new crank. Tocompensate for distortion in the design phase,the crankshaft was rough machined, heattreated and then finish machined. Thissignificantly reduced the concern, while finalmachining actually increased the strength ofthe component.

The ADI crank out performed the ductileiron version, showing no signs of fatiguecracking during bench testing. ADI’s internaldamping characteristics also gave the enginesuperior noise properties compared to engineswith steel or ductile iron crankshafts.Mechanical results performed on all threematerials are shown in table 2.

Independent truck stabiliser armTrucking in the Australian Outback is achallenging experience. The terrain is rough andisolated and distances may be exceptionallylong between service stops. When making the3,500km trip from Sydney to Perth, the priorityis to complete the journey safely while utilisingthe effective space in the trailer.

Originally, the independent suspensiondesign was a fabrication made from low carbonsteel. The first iteration was a 50mm thick, V-shaped swing arm. It was to be road testedover the route the truck would take in service -from Sydney to Perth and back. However, at thefirst trial of these components, the wheelssplayed under the truck during loading; thefully loaded trailer weighing 22.5 tonnes.

The test was run as planned but thewelded components failed after approximately1,200km. One of the suspension bracketsbegan to crack at its weld points; concern wasalso raised because these brackets would flexso heavily that the negative camber induceduneven tyre wear. A second set of welded steelbrackets was tested, travelling approximately4,000km before failure.

The suspension was re-designed as asingle-piece ductile iron casting; trial batcheswere cast and then austempered to achieveGrade 2 ADI (fig. 5). These parts weresubjected to the same Perth to Sydney trip asthe fabricated steel brackets. At the time ofwriting the ADI brackets had completed over322,000km without problems. As well asproviding an additional 20m3 storage spaceinside the truck, tyre life has been extended byover 80,000km.

Duralite truck hubWalther EMC has developed an ADI trucktrailer hub that is 2% lighter than itsaluminium counterpart and over 30% lower incost. This is a rather high visibility example ofADI replacing aluminium ‘kilo for kilo’, possiblybecause of ADI’s superior strength to weightratio.

Ford Mustang control armGM has demonstrated the feasibility of ADIsuspension control arms on Cadillaclimousines since 1995. More recently BentlerCorporation was contracted by Ford to producea lightweight, cost effective, independentsuspension system for its high performanceMustang Cobra sports car. ADI was chosen forthe upper control arms (fig. 6) for itscombination of low weight (approximately 3kg), noise damping and low manufacturingcost.

The components were FEA modelled totake full advantage of the strength and stiffnessof ADI. An aluminium design was consideredbut it was impossible to fit the much largerand thicker component in the given space. ADIallowed the manufacturers to meet their space,weight, safety critical, and cost objectives. Thesuccess of this application was recognised bythe 2003 AFS Casting Congress - at its annualCastings Contest, five of the top awards were

properties exceed those of forged, cast andmicroalloyed steels. Unlike aluminium,ADI’s endurance limit remains nearlyconstant after tens of millions of cycles.

• Improved noise damping: The presence ofgraphite in the ADI matrix improves noisedamping, for quieter, smoother runningcomponents.

• Superior wear and abrasion resistance:ADI’s abrasion resistance exceeds that ofconventionally processed steels and irons ata lower ‘bulk’ hardness level. Unlikecarburised steel, which loses wearresistance as the carburised layer isremoved, ADI improves in service. Wearresistance is superior to steel at any givenhardness level, making it ideal for earthmoving and other high abrasionapplications.

Case studies in transportTVR crankshaftThe TVR Tuscan Speed Six (fig. 3) made itsdebut in 1999 for sale to the UK and Japan asa right hand drive only vehicle. Striking inappearance and power, the Tuscan Sixaccelerates from 0-96.5km/h (0-60 mph) in 4.2seconds and reaches a top speed of 289.6km/h(180 mph). To achieve performance, thedesigners combined aerodynamic styling and ahigh power to weight ratio. Using composite

Fig. 2 The ADI ausferrite micro-structure - a mix of acicularferrite and carbon stabilised austenite

Fig. 3 Tuscan Speed Six (courtesy TVR Engineering Ltd)

Fig. 4 ADI crankshaft for the Tuscan Speed Six - 29.5kg

56 FTJ March 2004

Transport

given to ADI parts, including Best in Class forthe new Mustang suspension arm.

Another Best in Class was awarded for adrive wheel, part of a construction and

landscape utility loader. Originally an84 piece steel assembly, the redesignedcomponent is a one piece casting at a15% weight reduction, 55% lower incost. Thirty minutes of assembly timehave been eliminated while the partexhibits superior wear, durability andappearance compared to the original.

Looking aheadADI is being applied increasingly by theautomotive industries as the cost andperformance benefits are recognised.Following historical growth rates,annual world production of the materialis expected to reach 300,000 tonnes by2010, with approximately two thirds inNorth America. Carbidic ADI gradeshave recently been introduced and areadding to the interest worldwide.

Companies like ADI TreatmentsLtd actively help to develop the marketsand technology in partnership withfoundries and their customers. Some ofthe case studies illustrated have been

the result of such co-operations, frequentlystarting with workshops organised on thecustomer’s site. European manufacturers, oftenstimulated by USA connections, are now

Fig. 5 Pair of trailer suspension arms (courtesy Steele andLincoln Foundry)

Steel Ductile iron ADI ASTM ADIGrade 1

Yieldstrength 738 538 827 550MPa

Tensilestrength 910 903 1083 850Mpa

Fatiguestrength 400 324 427 N/AMpa

Impactenergy 325 75 141 100Joules

Elong (%) 23.2 10.8 13.7 10

Hardness 226-266 262-277 300 269-321BHN

Table 2. Test results for the steel, ductile iron, and ADI crankshafts, andASTM 897-90 Standard ADI Grade 1 specifications

implementing their own volume applications:Contract austempering services are set toexpand to meet the demand. The future for ADIis limited only by an ability to fully exploit thisunique material.

ADI Treatments Ltd, Doranda Way,West Bromwich B71 4LE.Tel: +44 (0)121 525 0303. Fax: +44 (0)121 525 0404.e-mail: arronrimmer@adit-uk.com