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The Ductile Iron News

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To Promote the production and application of ductile iron castings Issue 2, 2005

FEATURES

•Environmental Embrittlement ofDuctile Irons - A Review ofAvailable Data

• Tim Eilers Receives the DISAnnual Award

• Average Metalcasting Hiring Wage$9.58 According to AFS Survey

PDF ARTICLES

•Ductile Iron for Heavy SectionWind Mill Castings: EuropeanExperience

• Micro-shrinkage in Ductile Iron/Mechanism & Solution

•Applications of ADI in HighStrength Thin Wall AutomotiveParts

•Phase Transformation and FatigueProperties of Alloyed andUnalloyed Austempered DuctileIrons

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Tim Eilers (left) of Francis & Nygren Foundry Receivesthe Ductile Iron Society's Annual Award from DIS

President Alan Druschitz

Cover Story

View Ductile Iron Related Publications

Located in Strongsville, Ohio, USA15400 Pearl Road, Suite 234; Strongsville,Ohio 44136 Billing Address: 2802 Fisher Road, Columbus, Ohio 43204 Phone (440) 665-3686; Fax (440) 878-0070email:[email protected]

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Environmental Embrittlement of Ductile Irons - AReview of Available Data

by Martin Gagné, Manager - Sorelmetal Technical ServicesRio Tinto Iron & Titaniumand Kathy L. Hayrynen, Technical DirectorApplied Process Inc.

1. INTRODUCTIONThe generic definition of embrittlement is given by Krauss(1) as the“Reduction in the normal ductility of a metal due to a physical orchemical change”. Most common examples of such a phenomenonare blue brittleness, hydrogen (in situ) embrittlement and temperbrittleness. While the phenomena listed above referred to changes,chemical or physical, occurring within the metal structure, thisdefinition can be extended to embrittlement resulting from changesgenerated by the interaction between the metal and its environment.These types of embrittlement are identified as hydrogenembrittlement (hydrogen originating from the environment), stress-corrosion cracking and liquid metal embrittlement.

Ductile Irons have been recently reported to be prone toenvironmentally assisted embrittlement. The objective of this articleis to review the data published on this subject(2, 3, 4, 5, 6), with aparticular emphasis on “Austempered Ductile Irons”.

2. REVIEW OF ENVIRONMENTALLY ASSISTEDEMBRITTLEMENT MECHANISMSAs will be seen in the following sections, it is difficult to pinpointwhat mechanism (or interaction of mechanisms) causes theembrittlement of Ductile Irons in liquid environment. A quickreview of the known mechanisms is presented(7) to allow the readerto evaluate the possible relationship of such phenomena with thereported data.

2.1 Hydrogen EmbrittlementEmbrittlement may occur in many metals in presence of a verysmall amount of hydrogen. Hydrogen may be introduced duringmelting and entrapped during solidification, may be picked-upduring solid state processing (heat treatment, plating, welding…) orintroduced by a cathodic reaction during corrosion. Hydrogen ispresent in the metal as monoatomic hydrogen due to thedissociation of molecular hydrogen by chemisorption at the surface(or as a result of oxidation reaction (corrosion)).

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As illustrated in Figure 1, hydrogen causes delayed fracture at astress level significantly lower than the strength of the metal. Thefracture process may be cleavage, intergranular or transgranular andno single fracture mode is characteristic of hydrogen embrittlement.There is no univocal mechanism for hydrogen embrittlement. Oneof them proposed that hydrogen in relatively small amount


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decreases the cohesive strength in regions of high stresses.

Fig. 1. Delayed Fracture Curve Caused by HydrogenEmbrittlement(7)

2.2 Stress-Corrosion CrackingStress-corrosion cracking (SCC) is the failure of an alloy from thecombined effects of a corrosive environment and a static tensilestress (applied or residual). The chemical environment causing SCCdoes not produce chemical corrosion of the alloy and the speciescausing SCC need not be present in large concentration. Theformation and rupture of a passive layer at the crack tip is animportant mechanism. Moreover, it is widely believed thatelectrochemical dissolution plays a major role in the crack initiationand propagation. There is a possibility of the adsorption ofdamaging ions that weaken the atomic bonding at the crack tip. Ifhydrogen is generated as a result of corrosion species, it is thenable to enter the metal, diffuse to the crack tip and cause crackpropagation. There is growing evidence of a link between SCC andhydrogen embrittlement.

Figure 2 shows a schematic of crack growth under constant stressin a corrosive environment. Note that the threshold stress is lowerthan yield strength and that the resulting fracture pattern may beintergranular, transgranular or a mixed mode.

Fig. 2. Schematic of the Three-Stages Environmentally AssistedCracking under Constant Load in an Aggressive Environment(7).

2.3 Liquid Metal EmbrittlementLiquid metal embrittlement (LME) occurs when a solid metalsurface is wetted by a lower melting point metal and results fromthe direct interaction of the liquid metal atoms with the highlystrained atoms at a crack tip. Adsorption of atoms from liquidmetal greatly reduces the surface energy and though this, thefracture stress is reduced. The prerequisite for the occurrence ofLME are: i) a good intimate contact or wetting between the surface



of the solid metal and the liquid metal; ii) an applied or residualstress; iii) some measure of plastic flow and some obstacle todislocation motion (plastic flow) at the solid-liquid interface. Otherfactors reported to promote LME are the presence of a sharp notch,a high strain rate, coarse grain size and temperature(8).

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3. ENVIRONMENTALLY ASSISTED EMBRITTLEMENT OFDUCTILE IRONSTable 1 lists the various Ductile Iron grades investigated forenvironmentally assisted embrittlement by various authors. Alsoincluded are unpublished data recently obtained by Hayrynen andBoeri(9). All information presented in the following sections weresourced from these references.

TABLE 1Ductile Iron Grades Investigated for Environmentally AssistedEmbrittlement

Grades Ref. 2 Ref. 3 Ref. 4 Ref. 5 Ref. 6 Ref. 9

Ferritic X X X X F + 10%

P X ? Pearlitic X X X X

Q&T X X X ADI 1 X X X X X

ADI 2 X X ADI 3 X X X

ADI 4 X X X

ADI 750(MADI) X

Figure 3 compares the stress-strain behavior of normal andembrittled materials. The material property mainly affected is theductility which is significantly reduced when embrittlement occurs.Fracture toughness behaves similarly to elongation and was alsoused to quantify the embrittlement phenomenon. Data using bothproperties will be used in the following sections.

Fig. 3. Typical Stress-Strain Behavior of Normal and EmbrittledMaterials(2).

Note that the different authors may have utilized differentexperimental parameters and/or investigated the effect of differentvariables on the embrittlement. The presented analysis tried to takethese differences into consideration. However, please refer to theoriginal papers if more details are needed.



3.1 ASTM-A536 Ductile Iron Grades

· Ferritic Ductile IronFigure 4 presents the results of the investigation carried out byDruschitz et al(6) on 65-45-12 ferritic Ductile Iron tested at slowstrain rate (1% per minute) at room temperature, in contact withvarious liquid environments. Under these test conditions, noembrittlement was detected. Three other studies(2, 3, 4) reportedsimilar results, either by measuring tensile properties or fracturetoughness. It has been clearly shown by Komatsu et al(2) that theoccurrence of residual pearlite (~10%) does not affect the behaviorof ferritic irons vis-à-vis the embrittling effect of liquids.

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Fig. 4. Effect of Liquid Environments on the Elongation of FerriticDuctile Iron (6) (strain rate 1% per minute).

· Pearlitic Ductile IronsThe data reported on the embrittlement of pearlitic Ductile Irons incontact with water is presented in Table 2.

TABLE 2Effect of Water Environment on the Elongation of Pearlitic DuctileIrons (low strain rate)

Reference 3 4 6Dry Condition ~3% ~5% 7.5Wet Condition ~1% ~2% 6.8

It is seen that embrittlement of pearlitic Ductile Iron is reported byShibutani et al (3) and Martinez et al (4), but not by Druschitz (6). Itis worth noting that Komatsu (2) also observed the embrittlementphenomenon when measuring fracture toughness parameters.

A review of the process route used to produce the pearlitic DuctileIrons tested revealed that Druschitz (6) tempered his specimens at565oC (1050oF) for 2 hours while the other investigators (3, 4)

tested them as normalized. The tempering is carried out to reducethe hardness of the iron as shown in literature (10) and isaccompanied by a higher ductility of the structure. Druschitzreports 7.5% elongation after tempering, a high value whencompared to normalized/non tempered or as-cast pearlitic castings.



The description of the microstructure by Druschitz (fine pearlite)does not allow to conclude on the microstructural changes takingplace during tempering that would explain the increase in ductility.However, it appears unlikely that significant decomposition ofpearlite to ferrite + graphite occurred at the tempering temperatureaccording to data reported in literature (11). However, pearlitespheroidization, which is the precursor step to pearlitedecomposition and results in a quasi-continuous ferrite network, isknown to occur in less than two hours at 600-620oC in 3.5% C, 2%Si Ductile Iron and may explain the different embrittlementbehavior reported by Druschitz. Spheroidization of pearlite duringtempering for two hours at 565oC is considered probable byMullins (13).

· Quenched and Tempered Ductile IronFigures 5 a) and b) present the results obtained by Druschitz (6) andKomatsu (2) on quenched and tempered Ductile Irons. WhileDruschitz did not report the embrittlement of Q&T Ductile Iron,whatever the fluid used, severe embrittlement (measured by fracturetoughness) was observed by Komatsu (2). Note that Shibutani didnot report embrittlement when carrying out tensile test, probablybecause of the brittle behavior of the Q&T specimens used (noplasticity).

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Fig. 5 a)

Fig. 5 b)

Fig. 5. Effect of Liquid Environments on the Properties ofQuenched and Tempered Ductile Iron

a) Elongation (Druschitz (6))b) Fracture toughness (Komatsu (2))



The heat treatments used in the two studies were, however,different, the major difference being the tempering treatment, Table3. The treatment used by Druschitz resulted in a temperedmartensite structure in which carbides are fully spheroidized in aferritic matrix while the one used by Komatsu probably partlyretained the original plate-like martensitic structure.

TABLE 3Tempering Treatments Used for Q&T Specimens

Reference Temperature Time

Druschitz(6) 620oC (1145oF) 2 hours

Komatsu (2) 400oC (750oF) 1 hour

3.2 Austempered Ductile IronsVarious grades of Austempered Ductile Iron were studied forenvironmentally assisted embrittlement. Table 4 identifies them andthe investigators.

TABLE 4ADI Materials Evaluated

References 2 3 4 5 6 7ASTM-897M- X X X X X

1 X X X ASTM-897M- X X X

2 X ASTM-897M- X X

3 ASTM-897M-

4 SAE-J-2477-

750

· ASTM-897M GradesAll ASTM-897M Austempered Ductile Iron grades (1 to 4) wereembrittled when tested, either for tensile properties or fracturetoughness. As examples, data reported by Martinez et al (4),Druschitz et al (6) and Hayrynen et al (7) are presented in Figures 6and 7 and in Table 5. Table 6 compares the various relativereduction in elongation measured by each author.

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Fig. 6. Effect of Liquid Environments on the Elongation of ASTM-897M Grade 1 (Druschitz et al (6))

Fig. 7. Effect of Water Environment on the Elongation of ASTM-897M Grades 2, 3 and 4 (Martinez et al (4)).

TABLE 6Effect of Water Environment on UTS and Elongation of ASTM

897M Grades 1 and 3(Hayrynen and Boeri (8))

Grade Property Dry Wet1 UTS (MPa) 1012 575

Elong. % 11.2 3.4

3 UTS (MPa) 1274 1015

Elong. (%) 8.3 4.0

TABLE 7Relative Reduction of Elongation for ASTM-897M ADI Materials

when in Contact with Water

Grade References

2 3 4 5 6 7

1 80% 75% - 75% 72% 70%

2 - 82% 80% 73% - -

3 - 72% 83% - 52%

4 - 70% - - -

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It is seen in Table 7 that the reduction of ductility, as measured bytensile elongation, is about 70 to 80 %. The only data differingfrom the others is the one reported by Hayrynen and Boeri for ADIgrade 3. However, the initial elongation value was 8%, while datareported by Martinez et al mention about 11% as “non-embrittledvalue”. The rupture in wet environment occurred at the same value,i.e. 4%, in both studies.

It is interesting to note that all grades fractured at 3-4% elongationin all studies, with only a few exceptions. Finally, it is also worthnoting that the embrittlement by fluids other than water reported byDruschitz (6) was also reported by Martinez et al (5).

· SAE-J-2477-750 GradeThis grade of ADI, also referred as machinable ADI, wasinvestigated by Druschitz et al (6) and Hayrynen and Boeri (7).Druschitz investigated SAE-J-2477-750 materials with differenthardnesses; for comparison purpose with Hayrynen’s data, the dataon the more ductile material (247 BHN) is presented in Figure 8,while Hayrynen’s results are listed in Table 8.

Fig. 8. Effect of Fluid Environments on the Elongation of SAE-J-2477-750 ADI (Druschitz et al (6))

TABLE 8Effect of Water Environment on the Tensile Properties of SAE-J-

2477-750 ADI (Hayrynen and Boeri (7))

Sample Property Dry Wet1 UTS (MPa) 768 722

Elong. (%) 23.2 8.3

2 UTS (MPa) 698 659

Elong. (%) 22.7 7.9

When in contact with water, both studies report embrittlement ofthe material. However, as shown in Table 9, the “embrittled” valuemeasured by Druschitz is higher and the relative loss of ductilitysignificantly less than those reported by Hayrynen. Such adifference is believed to be related to differences inmicrostructures. These irons being austenitized in the inter-criticaltemperature zone, the matrix obtained prior to austempering is amixture of ferrite and austenite, the ferritic phase being not affected



by the austempering reaction. According to Druschitz (12), pro-eutectoid ferrite appeared mainly as a network in the final structureof its specimen, while it was more fragmented in the Hayrynen’ssamples (9). This may explain the different results reported by thetwo authors.

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TABLE 9Comparison of Data Reported on the J-2477-750 ADI Grade

Reference Druschitz et al (6) Hayrynen & Boeri (9)

"Dry" Ductility 18.5% 23.2% (*)

"Wet" Ductility 11.5% 8.1% (*)

Relative Reduction inDuctility 38% 65%

(*) Average of two tests

It is worth noting that Druschitz (6) reported no embrittlement whenin contact with liquids other than water base fluids.

4. DISCUSSIONThe review of the data available on the environmentally inducedembrittlement of Ductile Irons clearly shows that all grades, withthe exception of ferritic irons, are sensitive to the phenomenon. Themechanisms by which this phenomenon occurs are, however, notclearly identified. Nevertheless, the following observations fromthe published papers, together with the previously reported data,may help to understand the phenomenon.

i. The phenomenon is not a grain boundary embrittlement orrelated to the presence of graphite. Ferritic Ductile Irons,which exhibit a large area of grains boundaries and thehighest fraction of graphite are not sensitive to thephenomenon.

ii. The phenomenon is independent of the time spent in contactwith the liquid and disappears when the liquid is removed, asshown by Komatsu (2) and Martinez (4).

iii. Increasing strain rate (2, 4, 6) reduces the extent of thephenomenon, which is time sensitive and probably involvestime controlled phenomena (diffusion, chemical reaction,adsorption …).

iv. The phenomenon is observed in structures with extendedphase interfaces (pearlite, tempered martensite, ausferrite).

v. Shibutani (3) reported that similar embrittlement occurs whentensile test are carried out in a hydrogen atmosphere.

vi. Forming a ferritic layer at the surface of the specimensprevents embrittlement (3), as does painting the component(3).

Two mechanisms for environmentally assisted embrittlement ofDuctile Irons have been proposed in the literature: Hydrogen



embrittlement (3) and Chemisorption (5, 6) . At this stage, there is noclear evidence that would favour one theory via-à-vis the other.However, whatever the mechanisms involved, the presence of ahigh density of phase interfaces appears unavoidable for thephenomenon to occur. These phase boundary zones are privilegedsites for hydrogen pick-up and diffusion or for adsorption of atomsor molecules, both phenomena known to cause embrittlement. It isworth noting, however, that chemisorption (or Liquid MetalEmbrittlement) is favored by high strain rate, while the oppositewas observe for embrittlement of Ductile Irons.

The presence of ferrite, either as a fully ferritic matrix, as a quasi-continuous phase in a composite matrix (spheroidized pearlite, fullytempered martensite) or as a continuous network (J-2477-750 (6))probably prevents the embrittlement phenomenon by deformingplastically under the applied stress, avoiding the formation andpropagation of cracks.

5. CONCLUDING REMARKSThe review of available data on the environmentally inducedembrittlement of Ductile Irons indicates that high strength DuctileIron grades (pearlitic, Q&T and ADI) are sensitive to thephenomenon. When these grades are treated in order to get acontinuous (or quasi-continuous) ferritic matrix (by tempering orinter critical austenitization of ADI), the phenomenon fades.

The review of the data also evidences that three factors have to bepresent to observe the environmentally assisted embrittlement ofhigh strength Ductile Irons:

i. Presence of a liquid in contact with the material;

ii. Applied stress approaching yield strength of the material;

iii. Low strain rate.

When designing a part, the “design safety factor” usually ensuresthat the component will not be loaded near the yield strength of thematerial. Nevertheless, when working in a liquid environment,additional precautions should be taken to ensure that the threeabove-mentioned factors are not found at the same time. Paintingof the components should also be considered.

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REFERENCES

1. George Krauss, Steels: Heat Treatment and ProcessingPrinciples, ASM International, Materials Park, 1989.

2. S. Komatsu, C.Q. Zou, S. Shibutani and Y. Tanaka,“Embrittlement Characteristics of Fracture Toughness inDuctile Iron by Contact with Water”, Int. J. Cast Metals Res.,vol. 11, 1999, pp. 539-544.

3. S. Shibutani, S. Kamatsu and Y. Tanaka, “Embrittlement ofAustempered Spheroidal Graphite Cast Iron by Contact withWater and Resulting Preventing Method”, Int. J. Cast MetalsRes., vol. 11, 1999, pp. 579-585.

4. R.A. Martinez, R.E. Boeri and J.A. Kora, “Embrittlement of



Austempered Ductile Iron Caused by Contact with Water andOther Liquids”, Int. J. Cast Metals Res., vol. 13, 2000, pp. 9-15.

5. R.A. Martinez, S.N. Simison and R.E. Boeri,“Environmentally Assisted Embrittlement of ADI – CurrentUnderstanding”, Proceedings of the World Conference onADI, Louisville, 2002, pp. 91-96.

6. A.P. Druschitz and D.J. tenPas, “Effect of Liquids on theTensile Properties of Ductile Iron”, SAE Paper no. 2004-01-0793, SAE Conference, Detroit, 2004.

7. G.E. Dieter, Mechanical Metallurgy, 3rd Edition, McGrawHill, New York, 1986.

8. A.K. Sinha, Ferrous Physical Metallurgy, Butterworths,Boston, 1989.

9. K. Hayrynen and R. Boeri, Unpublished results, May 2005.

10. B.V. Kovacs, “Heat Treatment” in Ductile Iron Handbook,American Foundrymen’s Society, Des Plaines, 1999.

11. M. Gagné, “The Combined Effect of Chromium andManganese on the Thermal Stability of Eutectic andEutectoid Cementite in Ductile Irons”, CanadianMetallurgical Quarterly, vol. 25, no. 3, 1986.

12. A. Druschitz, personal communication, June 2005.

13. J.D. Mullins, personal communication, July 2005.

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Tim Eilers of Francis & Nygren Foundry receivesthe Ductile Iron Society's

Annual AwardIn 2003, the Ductile Iron Societyestablished a "Member Service Committee"and asked that their mission be to increasethe membership of the Ductile IronSociety. Actually, a committee wasdesignated before a committee chairmanwas found. A member came forth andvolunteered to head up this committee and the results have beenspectacular. DIS has added 8 new Foundry members and 5 newAssociate members in the past 18 months. A great deal of time andeffort went into designating potential members, establishing a planto contact them and the awesome task of following up with all thecompanies that showed an initial interest. The DIS rewarded thateffort by naming committee chairman Tim Eilers of Francis &Nygren Foundry to receive the Ductile Iron Society's AnnualAward.






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Average Metalcasting Hiring Wage $9.58According to AFS Survey

1695 N. Penny LaneSchaumburg, IL 60173www.afsinc.orgAugust 2, 2005847/824-0181

The 2005 American Foundry Society Wages and Benefits Surveyhas been tabulated, and its results provide a valuable insight intothe labor costs at metalcasting facilities throughout North America.Reports from 168 facilities representing 21,466 employees in theU.S. were used in developing the report, which collected data onwage averages for 28 job classifications.

According to the survey, on a national level, the lowest hiring wagewas $6/hr., the highest hiring wage was $15.65/hr., and the averagehiring wage was $9.58. The average hourly earnings were $14.05.Additionally, 27% of the respondents provided health care benefitsthrough an HMO insurance plan while 68% offered a PPO plan.The remaining respondents offered a combination of the two or areself-insured. The survey also tabulated responses to questionsregarding holidays and vacations, retirement plans, life insurance,sick leave and pay differentials per shift.

AFS uses the annual survey as an information-gathering springboard to keeping its members informed on the labor trends in theindustry. The survey is broken down into 20 labor market areas aswell as by type and size of metalcasting facility and union status.

AFS also partnered with The Management Assoc. of Illinois toprovide a National Executive Compensation Survey. For moreinformation on this survey or the AFS Wages and Benefits Survey,visit www.afsinc.org or contact Lynn Smith, AFS, at 847/824-0181or [email protected].

Headquartered in Schaumburg, Ill., AFS is a not-for-profittechnical and management society that has existed since 1896 toprovide and promote knowledge and services that strengthen themetalcasting industry for the ultimate benefit of its customers andsociety.

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2003 Keith Millis Symposium on Ductile Cast Iron

Ductile Iron for Heavy Section Wind Mill Castings: A European Experience

H. Roedter*and M. Gagné**

Rio Tinto Iron & Titanium Inc. *Frankfurt, Germany, **Montréal, Canada

ABSTRACT During the five year period from 1996 to 2001, the capacity of the installed wind energy plants worldwide has been multiplied by four, increasing from 6 x 103 MW to 25 x 103 MW. This increase has mainly taken place in Europe where many countries have embraced this source of energy as part of the solution to dependence upon foreign energy sources (gas and oil), environmental problems, and unsolved risks of nuclear power. The European Ductile Iron industry fully participated in this development. When compared to steel castings, Ductile Iron offers weight reduction and can meet the mechanical requirements of many wind mill parts (tensile, impact, fatigue) without heat treatment. Indeed, low Mn low P low Si ferritic Ductile Iron, produced under carefully controlled parameters, exceeds the targeted properties in the as-cast condition. The objectives of this paper are to review the development history of the Ductile Iron wind mill industry in Europe, to present typical wind mill applications for Ductile Iron and to describe the process parameters that allow the production of these castings in the as-cast condition. INTRODUCTION Since the beginning of history, humanity has explored a large number of sources of energy to assist in making every day work easier and more productive; one of the first sources used was wind. Indeed, Chinese built wind mills more than 4 000 years ago! However, the use of wind energy for electrical generation dates back to the 1880's when a Danish inventor converted classical wind mills to DC electricity generation. Today, more than 120 years later, the Danish wind turbine industry has a 50% share of the world market. The development of wind energy in Europe has been promoted by a number of factors. They include: • Reduction of dependence on oil producers; • Need for alternative sources of electricity to replace

coal, oil or gas, whose availability could be limited in the future and price subject to significant increase;

• Elimination of more risky technology such as nuclear energy;

• Free, renewable fuel; • Environmentally clean, although not completely problem

free; • Absence of waste generation. These factors, and others, have created conditions particularly favorable to the development of wind energy in Europe. The objectives of this paper are to present an outline of the development of the wind energy industry in Europe, to describe the major role played by Ductile Iron as a structural material in wind mills, to present a perspective of development of this industry in North America, and how it could benefit Ductile Iron foundries. DEVELOPMENT OF WIND POWER INDUSTRY IN EUROPE As previously indicated, wind energy for electricity generation dates back to the 1880's when Danish physicist Poul La Cour (1846-1908) converted classical wind mills to DC electricity generation. In the 1910's, wind mills for electricity generation became part of the Danish scenery with several hundreds operating wind mills generating 5 to 25 kW each. Construction of wind mills also took place in Germany and other European countries, but cheaper energy sources at an assumed nearly unlimited availability made the development of wind mill technology very slow. In the 1970's the oil crises (1973 and 1978), the increased concern of the public in general about the quality of the environment (global warming, acid rains,….) and the unsolved risks of nuclear energy forced the West-European governments to start looking at clean renewable energy. Nevertheless, it was 20 years after the first oil crisis that wind energy was clearly identified as an alternative source of energy to fossil and nuclear energy and, in the 1990's, construction of wind power plants boomed. Figure 1(1) presents the evolution of the installed wind energy capacity in Europe from 1997 to 2002; during that time period, the yearly installed capacity increased at an average rate of ~ 640 MW. A total of 32 037 MW, including the capacity installed before 1997, is in production in Europe. For comparison purposes, the installed capacity in North America is also shown in Figure


1; an increasing trend is seen during that period, but at a much lower rate. These developments led to the production of 16% of the total energy produced by wind in Denmark. In Germany, 2002 was a record year with 2 328 new units installed for a capacity of 3 247 MW; 12 000 MW of capacity is now in production representing 4.7% of the total electric energy produced. The turnover of this industry in Germany exceeds 3.5 billion US dollars.

Figure 1 : New Installed Wind Power Capacity in Europe and North America for the 1997-2002 Period. Projections for the period 2003-2007 show an acceleration of the installed capacity in Europe, as seen in Table 1, to reach a total approaching 60 000 MW in 2007.

Table 1 Projected Wind Mill Power Development 2003-2007 (1)

Year New Capacity Cumulative Annual Increase

MW MW % 2002 - 23 832 - 2003 6200 30 032 26.0 2004 6420 36 452 21.4 2005 6600 43 052 18.1 2006 7500 50 552 17.4 2007 8080 58 632 16.0

WIND POWER GENERATION: A MARKET FOR DUCTILE IRON The increase in demand for wind mills has stimulated the research and development in this area. Table 2 illustrates the evolution in size and capacity of the equipment, which went from 3.0 kW in 1980 to 1.5 MW in 2000 with experimental units of 4.5 MW to be put in service in 2005. Such a 4.5 MW unit is shown in Figure 2.

Table 2 Evolution of Wind Mill Characteristics

Year* Power Height Rotor Diameter

MW m (ft) m (ft) 1980 0.030 30 (98) 15 (49) 1985 0.080 40 (131) 20 (66) 1990 0.25 50 (164) 30 (98) 1995 0.60 78 (256) 46 (151) 2000 1.5 100 (328) 70 (230) 2005 4.5 124 (407) 114 (374)

* Estimated. Figure 2 : View of a 4.5 MW Prototype Unit (2). The increase in power and size has also resulted in the optimization of the structural materials used in the fabrication of the wind mills. It became rapidly clear that lighter weight materials meeting the required mechanical performance would be needed. Ductile Iron which offers a 10% weight reduction compared to steel was then considered by the designers for many components. Figure 3 presents a schematic of a wind power unit with a gear box. Gearless machines are also available and represent about 20% of the world market.

Figure 3 : Cross Section of a Wind Mill Unit with a Gear Box.

0

1000

2000

3000

4000

5000

6000

1996 1998 2000 2002 2004

Year

Inst

alle

d M

W p

er Y

ear Europe

North America


Table 3 lists the major Ductile Iron components utilized in the construction of wind power plants and the average weight of the parts per MW. Note that certain wind mill designs contain significantly more Ductile Iron than others. For example, the nacelle frame may weight up to 10 tons, which results in 19t of Ductile Iron castings in this particular 1 MW design. Note, however, that this ratio is usually not valid for large units (4.5 MW) for which bigger castings are required. All parts must meet the EN-GJS-400-18 LT specification, the impact toughness requirement being 12 J (8.8 ft-lb) (Charpy V-notch) at -20°C (-4oF). Examples of wind mill Ductile Iron castings are shown in Figures 4 (hubs) and 5 (shaft). Figure 4 : Ductile Iron Hubs for Wind Mills. Figure 5 : Ductile Iron Shaft for Wind Mills (390 kg).

Table 3 Ductile Iron Wind Mill Castings

Assembly D.I. Parts Average Weight

t/MW Rotor System Hub, blade

adapter, bearing 6.44

(hub : 4.5 t/MW) Shaft Shaft, bearings 0.44 Turbine Frame Nacelle, bed plate,

yaw ring 3.85

(nacelle : up to 10t)

Gear Box Housing, support, bearings

1.71

Others 1.00 TOTAL 13.44

MANUFACTURE OF WIND MILL CASTINGS Specifications and Properties Table 4 lists the minimum mechanical properties required for wind mill castings, in order to comply with the European EN-GJS-400-18LT (previously GGG 40.3) specification for ferritic Ductile Iron.

Table 4 EN-GJS-400-18 LT Specification for Wind Mill Castings Property Metric Imperial Tensile Strength 400 MPa 58 000 psi Yield Strength 240 MPa 35 000 psi Elongation 18% 18% Hardness 160-170 BHN 160-170 BHN Impact Energy at -20°C (-4°F) Average of 3 tests 12 J 8.8 ft-lb Minimum (1 specimen) 9 J 6.6 ft-lb The high ductility and toughness (as measured by V-notched impact Charpy tests) are of paramount importance for these castings because of the harsh weather conditions to which they may be exposed in Northern countries or offshore. In addition, the soft ferritic matrix, in conjunction with the lubricating effect of the graphite spheroids, eases the finishing machining operations to which the castings are submitted. In order to achieve the mechanical properties required for the castings, the microstructural targets are as follows: • Nodule Count: 100-200 nodules/mm2; excessively high

nodule counts impair impact strength, while very low nodule count may result in intercellular brittle phases that are also detrimental to strength and ductility;

• Nodularity: a minimum nodularity level of 95% is critical; poorly shaped nodules act as stress risers and as initiation sites for fracture under impact;

• Matrix: fully ferritic, without cell boundary carbides or pearlite due to segregation of trace elements, non-metallic inclusions or other constituents;

• No microshrinkage, or dross can be tolerated. Experience has shown that these microstructural characteristics can be achieved in the as-cast condition, avoiding a costly annealing heat treatment, by tightly controlling the chemical composition of the castings, by carefully selecting the charge materials and by optimizing the process parameters. Chemical composition Table 5 lists the typical chemical composition of finished wind energy castings.


Table 5 Chemical Composition of Wind Mill Castings

Element wt %

C 3.3 – 3.5 S 0.008 – 0.012 Si 1.9 – 2.2 P < 0.030

Mn < 0.15 Mg 0.040

The rationale behind the selection of this chemical composition is as follows: • Carbon: As shown in Figure 6, a high carbon content

increases the graphite fraction in the casting and lowers the absorbed impact energy, although lowering the transition temperature. A low carbon content is favorable to the upper shelf impact energy, but increases the transition temperature. The 3.3 – 3.5% C range was found to offer the best compromise to both properties.

• Silicon: The effect of silicon content on absorbed impact energy of ferritic Ductile Iron is shown in Figure 7. Because of the strong solid-solution hardening effect of silicon in ferrite, the lower the silicon content, the lower the transition temperature. However, a minimum silicon, although kept at a minimum level, is needed in the base iron to maintain its graphitization potential and as post-inoculant addition.

• Sulphur: It is recognized that a minimum sulphur content is required for inoculation (0.006%), while a sulphur content exceeding 0.012% may affect graphite nodularity. As shown in Figure 8, high nodularity is needed for optimum impact properties.

• Phosphorus: As shown in Figure 9, phosphorus is an embrittling element and every effort should be made to keep it to an absolute minimum. Phosphorus forms steadite, a low melting point eutectic, that segregates to cell boundaries during solidification. In heavy section castings, such as wind mill castings, the cell boundary regions can contain as much as 10 times the bulk P concentration. It is a requirement for such castings to maintain P content to less than 0.030% preferably below 0.025%. Although steadite may not be observed in the casting, phosphorus atoms present at cell boundaries reduce the cohesion level of the grain boundaries (4).

• Manganese: It is well known that manganese segregates to cell boundaries, especially in heavy section castings requiring long solidification time, and promotes the formation of pearlite in these regions. As shown in Figure 10, pearlite embrittles Ductile Iron and, in order to avoid its formation, 0.15% Mn is recommended as the maximum concentration in wind energy castings. Note that all other pearlite promoting

elements should be maintained at trace levels as well. The beneficial effect of low phosphorus content is also shown on this figure.

• Magnesium: Usually, a Mg content below 0.045% is recommended as this will ensure good nodularity while minimizing the formation of dross.

Figure 6 : Effect of Carbon Content on V-Notched Charpy Impact Energy of Ferritic Ductile Iron (3). Figure 7 : Effect of Silicon Content on V-Notched Charpy Impact Energy of Ferritic Ductile Iron (3). Figure 8 : Effect of Graphite Nodularity on V-Notched Charpy Impact Energy of Ferritic Ductile Iron (3).


Figure 9 : Effect of Phosphorus on V-Notched Charpy Impact Energy of Ferritic Ductile Irons (5). Figure 10 : Effect of Matrix Microstructure on V-Notched Charpy Impact Energy of Ductile Iron (3). Manufacturing Process The manufacturing process for as-cast wind energy castings does not differ significantly from the one used in foundries producing high quality ferritic Ductile Iron castings. Nevertheless, certain parameters need special attention in order to obtain, casting after casting, the same outstanding quality. Repeating and documenting every process step is key in succeeding in the manufacture of these demanding castings. • Charge Materials: A typical charge includes a minimum

of 40% high purity iron exhibiting consistently low concentrations of residual elements (P, Mn, Cu, Cr,…). Steel scrap, which is needed to lower the carbon content of the melt, should be certified low P low Mn steel, free of protective coatings (galvanizing, paint,….) and tramp elements. Ductile Iron returns from the production of wind energy castings are high quality materials that can be recycled in the charge. Preconditioning with SiC(6) is also recommended.

• Melting: In order to retain the required high metallurgical quality and nucleation potential of the liquid iron, superheating should be restricted to 1500°C (2730oF) for a short period of time.

• Magnesium Treatment: Treatment methods using elemental magnesium or high magnesium alloys can result in excessive deoxidation of the iron and are not recommended. Tundish cover ladle treatment that minimizes the amount of Mg alloy used is particularly well adapted for this type of production. One could consider the use of nickel-magnesium alloy. It has the advantage of minimizing the amount of silicon added at this process step while alloying with nickel is beneficial to impact energy without, at this level, enhancing pearlite formation.

• Inoculation: This is the key process step to ensure a high nodule count that would prevent intercellular segregation of phosphorus and pearlite promoting elements such as manganese, and avoid the formation of porosity. Powerful inoculants with good fading resistance should be used. Inoculation should be carried out in two steps: 0.4 to up to 0.8% inoculant during transfer to the pouring ladle and late inoculation (in-stream or in-mold) with 0.10 to 0.20% inoculant.

• Pouring Temperature: The castings are usually poured at 1370-1400°C (2500-2550°F), but the pouring temperature needs to be determined carefully, depending on the casting section size and risering technique used; the foundryman should take advantage of the self-feeding characteristics of a high metallurgical quality liquid Ductile Iron.

• Molding: Rigid, well compacted molds are key elements for the production of sound heavy section castings. Filters are often utilized to ensure inclusion free castings. Use of chills in critical areas and of mold wash is often practiced. Venting should be sufficient to ensure fast filling of the mold.

Figure 11 : Recommended Shake-Out Times as a Function of the Targeted Matrix and Modulus (7). • Shake-out: Cooling time in the mold is critical to

avoid/minimize pearlite. Figure 11 presents the recommended minimum shake-out time to obtain a

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Casting Modulus, cm

Ferrite

Pearlite

Solidification


ferritic structure as a function of the casting modulus. For comparison purpose, those required for a pearlitic matrix are also shown.

Repeating procedures and documenting process parameter data are essential components of the production of wind energy castings. Consistent, high quality castings can only be achieved by controlling each process step. The foundry must then be able to measure accurately, heat after heat, all the input materials, to control the process variables and to verify the final quality of the castings: • Base iron: Weight and chemistry controls of charge

materials, chemical composition within specified target, exact weight of the metal to be treated, close temperature control to meet the targeted pouring temperature range;

• Mg Treatment: Control of the alloy (chemistry, size,…), accurate weight control of the added amount;

• Inoculation: Control of the alloy (chemistry, size, …), exact and timely additions of the inoculant;

• Casting Quality Control: Mechanical testing and metallographic examination of test bars, hardness check of the castings.

PERSPECTIVE OF THE NORTH AMERICAN MARKET As shown in Figure 12, in 2002, USA produced about 15% of the world wind energy, which corresponds to 4 685 MW, or enough to serve 1.2 million households.

Figure 12 : World Wind Energy Generated by Region in 2002 (8). Like elsewhere in the world, wind energy in USA is largely dependent on government incentives. The lull in the US market is primarily due to the uncertain status of the wind energy tax credit (PTC), a key federal incentive. PTC was established in 1992 but expired twice during the past five years. The industry is seeking a multi-year extension that would provide a more stable investment environment in the US wind power market. Although this situation creates uncertainty, the American Wind Energy Association

(AWEA) forecasts positive growth of the installed capacity for the coming years. Figure 13 presents the AWEA forecast of the installed capacity for the period 2003-2010. In 2020, it is estimated that 6% of the US electricity will be sourced from wind power plants.

Figure 13 : Wind Energy Installed Capacity Forecast in North America (8). (Data for years 2008 and beyond is based on hypotheses that may change). Based on these previsions, forecasts have been made for three different Ductile Iron parts currently used in wind mill construction: hub, rotor flange and nacelle frame. Note that the number of units varies from one part to the other because certain wind mill designs use other materials than Ductile Iron in this application; that is particularly the case for the nacelle frame. Figure 14 presents the forecasts for Ductile Iron hubs, rotor flanges and nacelle frames needed in North America up to 2010. It is seen that, in 2010, more than 36 000 tons of Ductile Iron castings will be needed for these three parts. The total tonnage for all Ductile Iron components will exceed 50 000 tons.

Figure 14 : Projected Ductile Iron Needs for Specific Castings for the Period 2003-2012. The scenario presented in Figures 13 and 14 may be influenced by many factors, including the uncertainty regarding the PTC. The data presented assume that PTC

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will be renewed or, globally speaking, that the support of wind energy by North American governments will not drastically be reduced in the coming years. This is considered, from the PTC point of view, as a minimal scenario which could be positively influenced by new political measures (eg. reduction of CO2 emissions). Ductile Iron has a well established reputation in the wind energy casting market, due to the high quality castings delivered up to now. This performance and the relatively low cost of the material have contributed to protect the market from other materials. Maintenance of this quality level would confirm the dominating position of Ductile Iron in the future. CONCLUDING REMARKS • The development of wind power in Europe has accelerated, especially in the 1990’s, and this trend is expected to continue in the coming decade. Although the relative increase curve tends to flatten, the new installed power capacity is believed to exceed 9000 MW per year for many years to come. • Ductile Iron has been selected by the producers of wind energy power plants for many mechanical components because of its lower density and excellent reliability, even at low temperatures (-20oC) (-4°F). • For such applications, ferritic Ductile Iron with high impact toughness at -20oC (-4°F) (EN-GJS-400-18 LT) is the selected material for all Ductile Iron parts. • European foundries have shown that this grade can consistently be produced in the as-cast condition by carefully selecting charge materials, controlling the chemical composition and optimizing the manufacturing parameters. • The North American market for wind energy power plants is expected to follow the same trend as the European market during the 2000-2010 decade. • Ductile Iron will remain a material of choice for wind mill components in the North American market. In 2010, the tonnage of Ductile Iron hub, rotor flange and nacelle frame castings needed in the North American wind energy sector will exceed 36 000 tons. The total market for wind energy castings may then exceed 50 000 tons. REFERENCES 1. International Wind Energy Development & World

Market Update 2002 (Forecast 2003-2007) BTM Consult, March 2003.

2. Enercon E-112 aps Publication. 3. Ductile Iron Data for Design Engineers, Rio Tinto Iron &

Titanium, Montréal, 1998.

4. G. Kraus, Steels : Heat Treatment and Processing Principles, ASM International, Materials Park, OH., 1989.

5. Suggestions for Ductile Iron Production #68, Rio Tinto Iron & Titanium Inc., Montréal.

6. B.C. Godsell, AFS Transactions, vol. 82, 1978. 7. The Sorelmetal Book of Ductile Iron, Rio Tinto Iron &

Titanium, Montréal, 2003 (in print). 8. American Wind Energy Association Report, 2003.


Micro-shrinkage in Ductile Iron / Mechanism & Solution

R Siclari, T Margaria, E Berthelet, J Fourmann. PEM (Pechiney Electrometallurgie) – Paris, France

ABSTRACT The demand for increasing levels of quality in

iron castings continues relentlessly. Safety critical castings for the automotive

industry have two requirements : The first is linked to the qualitative aspects that bring the metallurgical performance, the second is purely economical; the cost of the casting directly influencing its survival.

Improvements of the mechanical properties of

the castings cannot be achieved by ignoring the influence of certain mechanical defects.

The need for control of chemical analysis and

microstructure is well documented and there are new tools available to the increasingly sophisticated foundry industry.

When considering the "health" of castings ;

micro-shrinkage, grain boundary dislocations, intergranular segregation, risk of inclusions etc... the state of knowledge of the cause of them and solutions to eliminate them, is quite weak. We would like to share with you , the advancements made in both research and industrial applications on the theme of the elimination of micro-shrinkage.

The work is supported by several industrial

examples. BACKGROUND

It has been admitted for some years that lanthanum has particular effects on cast iron compared to all other rare earth elements available. A large bibliography is at our disposal on this topic and more than 10 articles referred to the effects of pure lanthanum on iron properties, some of them are listed in the references attached. 1) Uses of Lanthanum 1-1 History

The use of lanthanum has been recognized since a long time to improve various iron properties. For example, a 1956 patent (ref.1) claims that the use of lanthanum alone (as for Sn, Pr or Nd without Ce)

improves machinability. Digging further into the topic, its main obvious metallurgical effects can be noted during nodularization and inoculation practices. Indeed, later in 1978 (ref.2 & 3) the use of lanthanum is identified as an efficient element to counter micro-shrinkage appearance. Since this discovery, the use of lanthanum has known a growing industrial development. 1-2 Practical uses

A typical example is given by the use of SIMAF, a pure lanthanum bearing FeSiMg used for in-mold nodularisation, patented by PEM in 1978 . Since this date, the use of SIMAF has developed to reach a consumption of 4000 t/y, corresponding to a SG iron production of 500 Kt/year of automotive safety parts in Europe. 1-3 Possible further uses of lanthanum element

The use of lanthanum bearing FeSiMg added up stream (ex: FeSiMg sandwich process) within the metallurgical foundry process has been explored. Laboratory experiments have shown: - rapid fade of the lanthanum content within molten iron, accentuated by the increase of iron temperature needed to accept the heat loss noted during conventional FeSiMg treatments. - difficulty in maintaining into an industrial environment the adequate and repeatable amount of lanthanum content in order to exercise on the molten iron the desire effects. 2) Behavior of Lanthanum (fading)

To illustrate the difficulty of controlling the fading property of the lanthanum introduced up-stream within the iron process, we must refer to the relative decrease of magnesium and lanthanum contents as shown on figure 1. The fading tendency of both elements have been studied during laboratory experiments, the curves are the following :


Figure 1: Relative concentration (Cr) decreases more rapidly for RE than for Mg - Cr is C(t)/Ci

C(t): concentration at a given time Ci: Initial concentration before fading

It can be noticed that an increase in residual Mg content will promote shrinkage formation (ref 4). 3) Late uses of Lanthanum

In order to solve this predominant fading tendency, the use of lanthanum has naturally found its applications during the iron treatments achieved at the latest stages of the metallurgical foundry processes.

Late additions of inoculants are nowadays well established in all foundry processes. Dedicated commercial products are also available for precise control of their uses. The effect of lanthanum on the iron solidification behavior and the optimization of its benefits can then easily be monitored and proven.

4) Metallurgical effects of Lanthanum 4.1 Solidification behavior

Through molten iron solidification process, different stages of grain growth occur:

Figure 2: solidification process of molten iron

First solids appear on mold walls, near locations of severe heat loss. Grains grow toward the opposite direction of the heat flow, usually perpendicular to mold walls. This growing behavior is named “columnar growth”, referring to the growing grain shape. Growth and definition of these grains are determined by the thermal properties of the mold.

When heat flow through mold walls has reduced, columnar growth is less favored. Particles of the columnar grains detach to grow freely in the liquid. This second solidification behavior is named “equiaxed growth”. Liquid iron convection motions and viscosity property dictate, in the mold, directions taken by growing grains. Within growing grain particles, heat flow and solidification directions are identical. Influencing factors: - high nucleation within molten iron favors equiaxed growth behavior as more solidification sites are available. - equiaxed growth behavior is favored by stirring motions (ref.5) - several other phenomena affect the transition point from columnar to equiaxed solidification behavior such as, heat flow characteristics, nature of solute elements at the molten iron solidifying front (segregation, …) This behavior is typical of primary solidification but applies also to the divorced Fe-C-Si eutectic (ref 6). 4.2 Experiments 4.2.1Principle Dimensional characterization of columnar solidification behavior growth zone at given solidification times and lanthanum addition rates will demonstrate the effect of lanthanum on iron solidification growth process. 4.2.2 Experimental design:

FeSiLa

Treatment and filling of the mold (2)

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Figure 3: Scheme of the experimental design An iron melt (1), is treated by 0,15% of late stream inoculant containing 0% up to 2,4 % of lanthanum (2). Treated melts are poured at the same temperature into cylindrical moulds. Solidification starts and progresses (3). At a given time then a given iron liquid/solid ratio, molds are emptied of their remaining semi-solid iron (4). Temperatures are monitored, specific solidification time is targeted to reach 50% iron liquid/solid ratio. Thicknesses of remaining solidified irons (shells) taken off from mold walls are measured. 4.2.3 Results

Resulting solid shells are compared:

Image 1: solid shell inoculated without La and its section view

Image 2: solid shell inoculated with La and its section view

Table 1: Thicknesses of resulting solid shells at 50% iron liquid/solid

ratio. (2 trials) 4.3 Conclusion

Lanthanum element promotes the equiaxed solidification behavior by: - developing a higher nucleation power, therefore more solidification sites within the molten iron. - modifying the molten iron viscosity favoring stirring motions within molten iron. - restricting the growth of columnar grains. This phenomenon is dictated by the physico-chemical properties of the Fe-La alloy binary diagram.

It has been demonstrated that alloying elements

showing a high m(k-1) * factor are able to slow down the grain growth rate . It is the case for Lanthanum. *m being the liquidus slope and k the distribution coefficient of the Fe-“element” binary diagram:

Figure 2: binary diagram Liquidus: T T m Cf= +

Solidus: T Tmk

Cf= +

Emptying (4) Solidification (3)


5) Consequences for shrinkage and micro-shrinkage At a given solidification stage, when lanthanum

is added, the equiaxed solidification behavior is favored. Consequences are that: - the thickness of the columnar solidification behavior grown zone is reduced, leaving larger free flowing passages for the remaining molten iron to travel within channels feeding casting areas in the one the need of liquid iron is essential to minimize the formation of shrinkage defects. - the semi-solid iron contains more solid particles. After having feed a given casting volume, less liquid is needed to compensate the resulting solid contraction. note: Thixotropic molding process used for the manufacturing of aluminum castings is based on same principals. 6) Others metallurgical effects It is known that Lanthanum has positive effect on nucleation (ref.7) and graphitization (ref.8). Lanthanum has the advantage of being an element able to generate stable sulphides and oxysulphides (ref 9 and 10) increasing the nucleation power of the liquid bath. 7) Industrial examples 7.1 Lanthanum bearing treatment alloy Lanthanum bearing inoculants are available under diverse chemical compositions and dedicated to the late stream inoculation practice (ex: FeSiLa, Win-4, etc.). 7.2 Examples of Lanthanum bearing alloy uses

7.2.1 Crankshafts ( courtesy Française de Mécanique)

Within the production of crankshafts, the effect of FeSiLa has been assessed. Results on microstructures and number of castings per mold (iron yield) are shown hereafter:

Image 3: SG iron structure without La showing micro-shrinkages

Image 4: SG iron structure with lanthanum clear of micro-shrinkages

Image 5: cluster of 4 crankshafts developed without lanthanum

Image 6: same crankshaft cluster redesigned after lanthanum addition. 6 crankshafts are placed within the same volume.


7.2.2 Further examples

Following are some images of castings for which FeSiLa had to be used to avoid micro-shrinkages.

Image 9 : Caliper

Image 10 : Steering Knuckle

Image 11 : Exhaust Manifold

Image 12 : Exhaust Manifold (flange view)

SUMMARY

Knowing and understanding lanthanum properties enable foundrymen to optimize their overall iron yields and to minimize the appearance of micro-shrinkages. The ideal use of lanthanum in late stage inoculation practices can be summarized as follow: - Lanthanum effect is consistant and more reproducible towards minimizing the risk of fading. - optimization of the inoculation treatment cost as the late stage inoculation rate would be smaller than all other upper stream inoculation practices that would have to take in consideration the dominant fading property of the lanthanum. - optimization of lanthanum nucleation and graphitization effects as they are naturally more effective when added late in the process. REFERENCES 1. patent US 2.970.902 – (1956)Production of ductile iron having improved machinability (using La based Rare Earth) 2. patent EP 4819 (1978) Ferrous alloys preparation with incorporation of Lanthanum for improved mechanical properties. 3. patent FR 024855 cast iron or steel containing small amount of Lanthanum improving their mechanical properties 4. R. Hummer, Cast Metals, volume2, number 2, 1988 5. Qiu P., Nomura H.and Takita M. Int. J. Casting Res., 2000, 13, 93-98 6. Ruxanda R., Beltran-Sanchez l. , . Masson J. and Stefanescu D.M., AFS Transactions 01-066, 1-12 7. Harding R.A.,Campbell J., Saunders N.J. The inoculation of ductile iron: a review of current


understanding. Conference; Solidification Processing Sheffied 97 4-6 july 8. Stefanescu D.M. , Loper Jr C.R. Effect of Lanthanum and Cerium on the structure of eutectic cast iron; American Foundrymen's society, Transactions, Vol 89, Cincinnati, Ohio Apr.-1 May 81 9. V Cochard. Inoculation od spheroidal graphite cast iron. Master of Philosophy. University of Birmigham. 1995 10. F Bachelot. Inoculation mechanisms of ductile iron. Master of Philosophy. University of Birmigham. 1997. 11. Zeji S.H. ,Yulin T., Binghuan J., Deshan L. The analysis of the effects of RE on graphite crystallization. International Symposium on new development of Cast Alloy Technology 12. Shen Z., Tang Y., Jia,g B., Li D., The vermiculizing effect of individuals Rare Earth elements La, Ce, Pr and Nd on graphite in cast iron conference on Rare Earth development and application vol 2 ADDITIONAL RESOURCES 1. Research works done at the IRC, Birmingham University in collaboration with PEM (J. Campbell, R Harding,F Bachelot, N Benailly, F Rossignol) 2. Internal research works done at the PEM Central Research Laboratory

2002 World Conference on ADI

Applications of ADI in high strength thin wall automotive partsR.A. Martínez, R.E. Boeri and J.A. Sikora

INTEMA - Faculty of EngineeringNational University of Mar del Plata – CONICET

Mar del Plata – ARGENTINA

ABSTRACT

The objective of this work is to discuss someconcepts referred to in the use of ADI for thin-wall, high-strength parts, and to summarize the development andapplication of two high strength automotive parts: hollowconnecting rods, used in a new two cylinder small carengine, and suspension front uprights, used ascomponents of racing car suspensions.

INTRODUCTION

The production of Austempered Ductile Iron (ADI)has shown a sustained rate of growth since its firstcommercial application in 1972. Currently, producers anddesigners continue to seek new applications, making useof the excellent characteristics of this emergingengineering material, that combines very high strengthwith good ductility, toughness, wear resistance andproduction advantages. A high number of successfulapplications in several fields, such as heavy truck,railroad, agricultural and automotive industries, as wellas new research results are regularly reported inconferences and journals, reflecting the technologicaland academic interest in ADI.

It is well known that a ductile iron part, which willbe heat-treated in order to obtain a good quality ADIcomponent, must be free from casting defects andcarbides, and have proper chemical composition,nodularity and nodule count.

In usual practice, ADI is mainly used to make partshaving thickness ranging approximately between 10 and100 mm. Its use in heavier section parts is restrictedbecause of the difficulty in obtaining the requiredaustemperability and the heterogeneous microstructuresobtained in slowly solidified heavily alloyed irons. On theother hand, the use of ADI in thin wall and high strengthparts (less than 4 mm) has not been reported yet, to thebest of our knowledge. This is probably due to thedifficulties to obtain sound, thin-wall pieces that are freefrom carbides. Carbides reduce the machinability in theas cast condition and deteriorate mechanical properties.Nevertheless, recent efforts have been aimed at theintroduction of ductile iron parts into the light partsmarket [1,2]. Light parts are mostly used in vehicles, andare customarily made of low-density metallic alloys orreinforced polymers. ADI is not a light metallic material,nevertheless, thin wall components can have strength to

density and elastic modulus to density ratios thatapproach those of cast aluminum, making it theoreticallypossible to apply ADI in high strength light parts.

The Metallurgy Division of INTEMA (DMI) hasbeen carrying out research in the field of ADI for severalyears. The research work of DMI is complemented bytechnical developments and technology transfer to theindustry.

The objectives of this work are to discuss aboutthe specific metallurgical issues that must be accountedfor before ADI is used in the fabrication of thin wall highstrength parts, and to summarize the experience gainedduring the development and application of two ADI parts.These parts are hollow connecting rods, used in a two-cylinder prototype engine, and front uprights, used ascomponents of racing car suspensions. Thedevelopment of both parts was required by the industrythrough agreements subscribed with the NationalUniversity of Mar del Plata, therefore, some data such aschemical compositions and details of heat treatmentcycles, were considered confidential and can not bereported in this paper.

METALLURGY OF THIN WALL ADI PARTS

The metallurgy of ADI has been extensivelystudied, and the general procedures to obtain a goodquality material are relatively well known. Nevertheless,when parts to be austempered are thin walled, someparticularities may have to be accounted for. In order toevaluate the feasibility of using ADI for high strength thinwall parts, the role of the following factors needs to bediscussed in advance:

a) Chemical compositionb) Molding and castingc) Solidification structure, microsegregation and

carbidesd) Solid state Transformation kinetics, final

microstructure and properties

Chemical composition:The chemical composition of the ductile iron

used to cast thin wall parts must be selected in order tosatisfy several requirements. Castability must bemaximized in order to assure complete mold filling. Tosatisfy this requirement it is advisable to usehypereutectic alloys (carbon equivalent: 4.4 - 4.6 %).


Meanwhile, it is also advisable to limit the silicon contentin order to avoid detrimental effects in the ductility.Regarding other alloying elements, the austemperabilityrequired to produce thin wall ADI is minimal, therefore,the use of alloying elements is unnecessary. Theconditioning of the melt may also be specific to thisapplication, in particular what relates to nodularizationand inoculation.

Molding and castingProper design of the mold and casting procedure

is critical when making thin wall ductile iron. Any internalor surface defect will have a strong detrimental effect in ahigh strength part. Dimensional tolerances are also muchmore severe. This requires the use of high precisionmolding methods, and careful design of the moldassembly, so as to guarantee proper positioning ofcores. Careful study of the metal flow during filling is alsorecommended. Some studies have evaluated the use ofcounter gravity casting procedures to improve flow duringfilling [3]. Different aspects must be taken into accountduring the fabrication process of the parts such as: molddesign, preheating of molds, use of heaters, filters andthe painting of the mold cavities to obtain good surfacefinish. Each type of part needs its specific analysis.

Solidification structure, microsegregation andcarbides:

It is well known that, as a result of the increasedsolidification rate imposed on the ductile iron when castin thin walled parts, the nodule count rises noticeably [1].Figure 1 shows the variation in nodule count for differentthickness of the same melt, cast in the metal castinglaboratory of DMI [4]. A 3mm thickness part can shownodule counts up to 1000 nodules per square millimeter,as compared to the 100-200 nodules per squaremillimeter usually found in conventional parts of about 25mm thickness.

The solidification structure is other feature thatmay be significantly different in thin wall ductile iron. Infact, recent studies showed that the Last To Freeze(LTF) melt locates at intra-dendritic liquid pools [5]. Thedispersion of the LTF becomes finer as the solidificationrate increases. The LTF shows higher concentration ofalloying elements of direct segregation behavior [6], andwill have greater tendency to show inclusions andmicrovoids. The austenite dendrites arm spacing isaffected by the solidification rate. The finer the austeniticmicrostructure, the higher the dispersion ofmicrosegregation. as it was confirmed by means of EDXdeterminations and qualitative color metallographytechniques [7]. These are positive effects of the highsolidification rate. On the other hand, there are alsonegative effects. The high solidification rate stronglypromotes carbide formation [8]. Carbides are detrimentalto the ductility, the toughness, and the machinability of

DI. However, since base material is unalloyed DI, thesecarbides will show small contents of carbide stabilizingelements, and are basically unalloyed cementite. Thissuggests that a dissolution annealing could easilyeliminate them. In addition, the large density of graphiteparticles, that act as carbon sinks during carbidedissolution, and the short distance between them, wouldalso favor dissolution during annealing [9].

FIGURE 1. Nodule count as a function of thickness

Solid state transformation kinetics, finalmicrostructure and properties:

The noticeable raise in nodule count causes anincrease of graphite-matrix interphase area. Graphite-matrix interphases provide heterogeneous nucleationsites. This may accelerate all phase transformationsinvolving solid-state nucleation, such as austeniteprecipitation during austenitization heating, as well asferrite precipitation during austempering. In addition, thepresence of larger nodule counts will decrease thedistance for carbon diffusion from the matrix to thenodules. The increased heterogeneous nucleation rateadded to the shorter diffusion distances, could lead to asignificant increment in the rate of solid-statetransformations.

In the specific case of austempering of thin wallDI, it is expected that the faster transformation kineticsand the smaller microsegregation present in the matrix,would contribute to eliminate unreacted austenite poolstypically present at the LTF, producing a finer and morehomogeneous final microstructure, and perhaps bettermechanical properties than those found in regularthickness DI. On the other hand, the austempering timeneeded to obtain optimum mechanical properties may bedifferent from that measured for DI of regularmicrostructure.

T hickness [ m m ]

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PROJECT FEASIBILITY

The analysis above shows a number of factorsinherent to thin wall ADI production that should beclarified before an application is foreseen. Nevertheless,the matters to be studied are certainly complex andwould require lengthy and careful investigations. Thetime available for the development of the parts, as usualin most cases, was limited and relatively short. On theother hand, most particular factors discussed earlierwould suggest that it is possible to obtain betterproperties in thin walled than in regular DI. Therefore, itwas decided to complete the development on the basisof the existing knowledge, making sure that the possibledetrimental effects are not affecting the partperformance.

DEVELOPMENT OF HOLLOW CONECTING RODS

The first part to be developed was a hollowconnecting rod (Figure 2), as required by the automotiveengineering enterprise PRONELLO I+D. This connecting

FIGURE 2. Longitudinal cut of the hollow connecting rod.

rod was a critical part of the design of an innovative twocylinder engine, that generates 55HP at 5500 rpm. Theuse of hollow connecting rods allowed the designengineer to maximize the moment of inertia of the part,while minimizing the alternative mass. These interestingfeatures required a castable, inexpensive high strengthmaterial. ADI was clearly the material of choice. Initialrequirements of tensile properties and fatigue strengthdefined by the design engineer could be satisfied by agrade 1 ADI (ASTM 897M-90). The thickness of mainlyportions of the part was as small as 3mm. Thisinnovative design allowed decreasing the weight of theconnecting rod from 600 grams, if made of massiveforged steel, to 400 grams.

The metallurgical development of the connectingrod was initiated under the following premises:

! The parts should have good machinability as cast.! The parts should show small and repetitive

dimensional variation after austempering.! Development of parts, including models, molds,

castings, tests, etc, should be completed in fourmonths.

! The mechanical properties of the parts should matchdesign values.

Additionally, the requirement for the parts to be freefrom casting defects is, in the present case, challenging,considering the complex part geometry.

FIGURE 3. Sand mold used to obtain the connectingrods.

Design, of models, molds, casting and heattreatment cycles

The first step was to achieve parts of ductile ironfree from defects. Models and molds were carefullydesigned and prepared. At this stage of the developmentan eutectic alloy was used. The melt was poured intoindividual molds to evaluate the influence of thesolidification rate on the microstructure of differentsections. Figure 3 shows the final sand moldconfiguration as was used to manufacture connectingrod prototypes. First castings showed pearlitic matriceswith abundant ledeburitic carbides, mostly located at the90-degree corners of the hollow sections. This firstexperiments showed the need to take measures todiminish carbide precipitation. Additionally, in order toreach a better control of the dimensional changesproduced during the ADI heat treatment, it was advisableto carry out the austempering treatment starting fromparts having mainly ferritic matrix, as suggested byMoncada et al [10].

A decrease in the presence of carbides and anincrease in the amount of as cast ferrite was obtained byincreasing the equivalent carbon (CE), as suggested by


the literature. C and Si contents were set to have an CE=4.4 – 4.5 %, while Mn content was decreased to theminimum levels compatible with industrial productionprocesses. No other alloying elements were added. Evenusing a hypereutectic alloy, no problems of graphiteflotation were evident since solidification takes placerapidly. Further changes included modification of themolds, in order to obtain a good filling and to diminish thecooling rate. The introduction of these modifications wasvery successful and the new set of parts resulted inmainly ferritic matrices, free from defects, with lowamount of carbides, and with nodule counts higher than700 nod/mm2 in the thinner sections.

Along with the parts, ½ inch “Y” blocks were castwith the same molten metal, and cylindrical tensilespecimens (1/4”ASTM A370-94), were machined. Thespecimens were heat treated together with thecomponents.

Austempering heat treatment cycles weredesigned to produce ASTM 897M 850-550-10 ADI. Firstcycles were applied to test specimens, and the resultingproperties were measured through tensile testing.Connecting rods were heat-treated using the samethermal cycle that gave satisfactory properties on thetensile tests.

Examination of cast prototypesFour sets of connecting rods were cast. Some of

the parts of the first set had to be rejected because ofimproper geometry, caused by the misalignment of thecores. Other small defects, such as pores and surfaceinclusions, where also found. These problems weresolved by introducing changes in the mold assembly, andthe remaining sets gave sound parts.

One part of each set was sectioned in severalsections to examine the as cast microstrusture. Matrixmicrostructure was mostly ferritic. Only small amounts ofcarbides were observed, preferentially located at the 90°corners of the rod hollow stem.

The machinability of as cast parts was alsoexamined and found to be satisfactory, since locations atwhich the rods need to be machined are relatively thickand showed no carbides.

Controls and Tests carried out after heat treatmentAustempering heat treatments cycles were applied

to the parts following the results on the tensile testspecimens.

Some rods were sectioned and examinedmicroscopically. The final matrix micostructure was fullyausferritic, and free from carbides at all locations. Thisproves that the austemperability is sufficient for theapplication, and confirms that the small amount ofcarbides present as cast, effectively dissolve duringaustenitizing. The microstructure was also almost free ofunreacted austenite pools, as shown in Figure 4. Figure

5 shows the ADI microstructure obtained on a tensile testsample taken from the ½ inch “Y” block, heat treatedtogether with the parts. Comparing Figures 4 and 5, it isclear that the microstructure of larger section thicknessparts shows unreacted austenite pools, and those arenot present in the finer microstructure of the thinner wallsamples. This also confirms the existence of differencein the solid state transformations as a result of thedifferent initial microstructure.

FIGURE 4. Microstructure of thin wall ADI (400x)

FIGURE 5. Microstructure of ½ inch “Y” block (400x)

The change of the length of the connecting rod asa result of the austempering was measured. The resultsshowed a negative change of approximately –0.1%. Thesmall change in dimensions allowed to perform almost allmachining before heat treatment. The only operationperformed after austempering was the grinding of thesurfaces of piston pin and crankshaft pin housings.

Assembled connecting rods (using Maraging steelbolts) were tested in tension using the set up shown in


Figure 6. Nine parts were tested. All fracture events tookplace on the bolts (Figure 7), at loads between 11.150and 11.500 Kg. Based on this results, and through themodeling by finite element codes, it was established thatfor these loads, the maximum stress at the rod reachedapproximately 870 MPa. This level of tensile strengthconfirmed that actual UTS of the part is high andcomparable to that measured on standardized tensiletest specimens.

Some connecting rods were assembled in theprototype engines, and were extensively tested in servicewith excellent results. The new engine was presented inseveral national and international expositions.

FIGURE 6. Assembly used for static testing.

FIGURE 7. Connecting rod after test showing a brokenbolt

DEVELOPMENT OF HOLLOW FRONT UP RIGHTS

The very good result of the development of theconnecting rods encouraged the designers to apply ADIin the construction of other complex geometry hollowhigh strength parts. Front uprights for racing cars weredeveloped. The parts were going to integrate a highstiffness front wheel suspension and steering assembly.They needed to be strong, light and reliable. As in thecase of the connecting rods, the development needed tobe completed in four months. Again, the great flexibilityof the casting and heat treatment processes was a keyissue favoring the use of ADI, since these allowed thefast introduction of changes in the design and propertiesof the parts.

Parts were designed based on the mechanicalproperties of an ADI Grade 2 (ASTM 897M-90). Thecalculations accounted for the very high loading on thispart as a result of inertial forces.

FIGURE 8. Front upright design

The steps of the development were similar tothose followed for the connecting rods. Figure 8 showsthe front uprights. The parts were tested statically toverify their capability to withstand the maximum designloads, and to measure their stiffness. Figure 9 shows theload–displacement curve for a specific loadconfiguration, which exhibits proportionality until 2800kg.This value surpassed the load demanded by thedesigner. Nevertheless, stiffness needed to beincreased. This led to a redesign process, which endedin a slightly heavier, thicker walled part.


FIGURE 9. Ad Hoc flexion test of front upright.

The part was successfully introduced in some ofthe racing cars of the Turismo de Carretera NationalCategory of Argentina during the 2001 championships.Figure 10 shows an upright assembled in a racing car.The performance of the cars was strongly improved bythe new design, and they obtained several first positions,pole positions and lap time records. The parts have beenredesigned and adapted to cars of several racing teamsrepresenting different brands of carmakers. The redesign

process continues.

FIGURE 10. Front upright assembled in the frontsuspension

FUTURE WORKIn spite of the relative lack of specific knowledge

about the metallurgy of thin wall DI, the twodevelopments of parts described above were successful.Nevertheless, this does not imply that furtherinvestigation on the metallurgy of DI is not necessary.Indeed, these successful applications encourage

carrying out specific research to make possible a morepredictable and generalized use of thin wall ADI. TheMetallurgy Division of INTEMA begun research aimed tostudy the transformation kinetics of high nodule count DI,and to examine the carbide dissolution rate in thismaterial.

SUMMARY AND CONCLUSIONSA hollow connecting rod for a two-cylinder car

engine and a front upright for a racing car were made ofthin wall ADI. The development was based on theexisting knowledge for regular ductile iron, which is notentirely applicable to this case.

The success in the development confirms theapplicability of ADI to build complex thin walled parts ofhigh strength. Design engineers should be aware of theexcellent properties of ADI, to be able to make use of therelative advantages of this material.

REFERENCES1- K. Davis, M. Sahoo, A. Javaid. “Effect of chemistry

and processing variables on the mechanicalproperties of thin wall iron castings” AFS CastingCongress (2000).

2- C. Labrecque, M. Gagné. “Development of carbidefree thin wall ductile iron castings” AFS CastingCongress (2000).

3- J. Massone, J. León Torres, L. Dix, D. Stefanescu.“Development of counter gravity casting for thin wallductile iron “. Proceedings of the InternationalConference on the Science of Casting andSolidification. Pp 353-360. Romania. 2001.

4- L. Achour, M. Martinez Gamba, R. Boeri, J. Sikora“Thin wall ductile cast iron; casting methodology andmicrostructure characterization” Proceedings ofIberomet Congress. Pp 419- 426. Spain. (200).

5- R. Boeri and J. Sikora “Solidification macrostructureof spheroidal graphite cast iron” Int. Journal of CastMetals Research, Vol 13 pp 307-313. (2001).

6- R. Boeri and J. Sikora. “Influence of austenite matrixmicrostructure on the phase transformations duringaustempering of ductile iron”. Proceedings of FirstInternational Conference on Retained Austenite”pp497-502. (2000).

7- G. Rivera, R. Boeri and J. Sikora. “ Influence of thesolidification microstructure on the mechanicalproperties of ductile iron” Int. Journal of Cast MetalsResearch, Vol 11 pp 533-538. (1999).

8- A. Javaid, J. Thompson, M. Sahoo, K. Davis.“Factors affecting the formation of carbides in thinwall ductile iron castings. MTL 98-24 (OP-J),CANMET (1998).

9- A. Giacopini, R. Boeri and J. Sikora. “Carbidedissolution in thin wall ductile iron” Unpublished.Submited to ISIJ International (2002).

10- O. Moncada and J. Sikora “Dimensional Change inADI” AFS Transactions. Vol. 104, pp 577-580 (1996).

DEFLECTION [mm]DEFLECTION [mm]DEFLECTION [mm]DEFLECTION [mm]

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2002 World Conference on ADI Paper Template

Phase Transformation and Fatigue Properties of Alloyed and Unalloyed Austempered Ductile Irons

by Chang-Zhin Wu and Teng-Shih Shih

ABSTRACT Specimens of ductile irons austempered at 320℃ and 360℃ for 2 hrs were conducted with ultrasonic vibration treatment and high cycle fatigue (HCF) testing. Experimental results show that specimens after ultrasonic vibration treatment progressively developed phase transformation induced by plastic deformation. Some stringer-type austenite was found to precipitate carbides, while some island-like austenite was transformed into martensite. The best HCF strength obtained in the alloyed specimens at 360℃ was attributed to its largest amounts of retained austenite with good toughness. This retained austenite was likely to transform to martensite under plastic deformation, resulting in a greater fatigue crack growth resistance. INTRODUCTION The austempering reaction of ADI usually occurs in two stages in time sequence 1. In Stage I, acicular ferrite (α) initially nucleates and grows in the vicinity of graphite and along the grain boundary of prior austenite. With the growth of ferrite, carbon diffuses to austenite and forms carbon-enriched stabilized austenite (γhc). In Stage II, carbon-enriched stabilized austenite (γhc) decomposes to ferrite (α) and carbide. Extending the protracted austempering time to Stage II produces a large fraction of ferrite and carbide precipitate.

It is believed that work hardening and stress-induced phase transformation in the matrix are in relation with cavitation erosion2-3. Wantang et al.4 demonstrated that structural changes occur after cavitation erosion for a Cr-Mn-Ni stainless steel. The weight fraction of martensite in the surface layer being exposed to cavitation erosion increases gradually.

Some research on high-cycle fatigue behavior of ADI indicated that the fatigue limit increased with the toughness and retained austenite content due to the high strain-hardening nature of austenite5-6. Aranzabal et al7 also remarked that large amounts of retained austenite would undergo martensitic transformation under plastic deformation in ADI austempered at higher temperature and with a greater content of retained austenite. The present research aims to adopt an ultrasonic vibration treatment to investigate the phase transformation in the matrix microstructure of ADI. HCF fatigue testing was also tested to obtain the fatigue stress-life (S-N) curves and to assess fatigue resistance of alloyed/unalloyed ADI. EXPERIMENTAL PROCEDURE The chemical compositions of ductile irons used in this study are given in Table 1. The ductile irons were

poured into 1-inch Y-block sand molds. All specimens were austenitized at 900℃ for 1.5 hrs and then subjected to isothermal heat treatment at 320℃ and 360℃ for 2 hrs. Table 1 Chemical compositions of ductile iron

Specimens C% Si% Mn% S% P% Cu% Ni% Mg% Alloyed 3.65 2.76 0.36 0.014 0.016 0.6 0.9 0.045

Unalloyed 3.51 2.45 0.30 0.010 0.017 - - 0.036

Groups of specimens were subjected to 46 KHz ultrasonic vibration at regular intervals of about 5 minutes. The morphologies and microstructural changes in the matrix were observed and assessed by using optical (OM) and/or scanning electron microscopy (SEM). Changes in retained austenite content were measured by X-ray diffraction using a Cu-target tube. A Vickers hardness tester at a test load of 25g was also employed to measure in areas near graphite and in intercellular regions. HCF specimens were tested in a rotary bending fatigue testing machine at a frequency of 2400 rpm (40 Hz) for high-cycle fatigue test. The test was run to failure or to 107 cycles at which specimens were considered to be a run-out. S-N curves were established for each specimen. Eight to ten specimens of each material condition were tested at sufficient stress levels to obtain reliable S-N curves and fatigue limits. RESULTS AND DISCUSSION Phase transformations during ultrasonic treatment Ultrasonic waves are periodic sound waves consisting of cycles of compression and expansion. Each cycle exerts a positive pressure and a negative pressure respectively. If a large negative pressure overcomes the liquid’s tensile strength at an expansion cycle, a cavity may form in the liquid8. When a cavitation bubble collapses at a location away from any solid boundaries, it does symmetrically and emits a shock wave into the surrounding liquid. When a cavitation bubble contacts with or is close to a solid surface, it will collapse asymmetrically and forms a micro-jet impacting directly toward the solid9.

Micro-jet impacts and shock waves cause stringer-type stabilized austenite, which was rich in carbon, to undergo Stage II phase transformation,γhc→α ＋ carbide. After ultrasonic treatment, the temperature of the water solution was slightly increased from 29℃ to about 40℃ for a treatment time of 10-15 minutes. Micro-jet impact and shock waves may locally raise the temperature of the matrix (ausferrite) and promote the diffusion of carbon to carry out Stage II transformation. Stringer-type austenite gradually

decomposed into ferrite and carbides. The micro-jet impact and shock wave energy forced

the island-like austenite to carry out a martensitic transformation. Fig. 1 shows the microstructure of ADI, which had been treated at an austempering temperature

of 360℃ for 2 hrs after ultrasonically treatment for 15 minutes. Most researchers agree that ADI exhibits a transformation-induced plasticity effect (TRIP), and martensite formed by deformation was mainly generated from the island-like austenite being present in the microstructure7,10. Because a lower carbon content existed in the interior of island-like austenite, less strain energy was necessary to produce the mechanical transformation of austenite and less energy had consumed the TRIP. High-cycle fatigue strength Results of the rotary bending fatigue S-N curves for the specimens are shown in Figs. 2a and 2b. The best-fit lines were used in all subsequent analysis of HCF properties. The estimated fatigue limits are also given in Table 2. The estimated fatigue limits at the high austempering temperature (360℃ ) were higher than those tested at the low temperature (320℃). In addition, the alloyed specimens obtained greater high cycle fatigue strength than the unalloyed specimens.

The fatigue properties are indeed affected by the matrix microstructure of ADI, especially the amount of retained austenite. At lower austempering temperatures, slim acicular ferrite plates are densely distributed in the matrix with a lower volume fraction of retained austenite. Contrarily, at higher austempering temperatures, the matrix presents a coarse ferrite lath and a larger volume fraction of retained austenite. Greater amounts of retained austenite created more barriers for fatigue crack growth and fatigue life was extended. This is also related to the plasticity-induced martensitic transformation of unstable retained austenite during HCF testing. And attributed to the addition of copper and nickel, the alloyed specimens presented a more homogeneous matrix microstructure with large amounts of retained austenite. The addition of alloys formed a diffusion barrier to hinder carbon diffusion at the graphite interface during solid-state transformation. It delayed Stage II reaction, thereby creating a wider processing window for the achievement of high toughness performance. As alloyed specimens austempered at 360℃ had much more blocky retained

10000 100000 1000000 10000000 100000000Cycle to Failure

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Pa)

500

Ni-Cu 360C

Unalloy 360C

Fig. 1 Microstructure of alloyed ADI 360C/2hafter ultrasonic treatment for 15 minutes

island-like carbide particles

Stringer type

Fig. 2a S-N curve of alloyed/unalloyed specimens at austempering temperature of 360C.

10000 100000 1000000 10000000 100000000Cycle to Failure

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Fig. 2b S-N curves of alloyed specimens at differentaustempering temperature 320℃ and 360℃

austenite, more stress-induced martensitic transformation occurred during HCF testing and led to a better HCF performance. Micro-hardness variations Variations of the micro-hardness values in ADI specimens with ultrasonic cumulative treatment time are shown in Fig. 3. The effect of ultrasonic micro-jet caused the alloyed specimens to present a slower and smoother increase of the micro-hardness values with increasing treating time, because alloyed ADI needs to absorb more energy to have carbide precipitates or stress-induced martensitic transformation. HCF testing also reveals that alloyed specimens have better fatigue strengths; that is, they will fracture only by absorbing more energy. So from the viewpoint of energy, Table 2 Estimated fatigue limits of alloyed/unalloyed

specimens austempered at 320℃ or 360℃.

Specimens Austempering

condition ℃

Volume fraction of retained

austenite(%)

Fatigue limit

(MPa)

360 27.7 350 Alloyed 320 18.2 270 360 21.6 290 Unalloyed 320 13.2 220

increasing austempering temperature and alloy (Ni and Cu) addition lead to an increase in retained austenite content. Retained austenite performs a great work hardening capability and can absorb energy resulting in better toughness and fatigue strength. SUMMARY 1. After subjecting samples to an ultrasonic treatment,

some island-like austenite is found to undergo a phase transformation of austenite to martensite induced by shear stresses from impact of micro-jets and shock waves, while some stringer-type austenite is found to precipitate carbides.

2. HCF testing shows that higher austempering temperature and alloy addition result in an increase in the retained austenite content and excellent toughness and fatigue strength.

3. After the process, micro-hardness values of alloyed ADI are slowly increased with increasing cumulative treatment time. It demonstrates that alloyed ADI needs more energy to cause phase transformation. This viewpoint corresponds to the result of HCF testing in which alloyed ADI also needs more energy to reach fracture.

REFERENCES 1. J. F. Janowak and R. B. Gundlach, “Development of a Ductile Iron for Commercial Austempering”, AFS

Transactions vol. 54, 1983, pp 377-378. 2. G. B. Jiang, Y. K. Zheng, Y. Y. Yang, and H. S. Fang, “Cavitation Erosion of Bainitic Steel”, Wear vol. 215, 1998, pp 46-53. 3. Marcio Gustavo Di Vernieri Cuppari et. al., “Correlation between Microstructure and Cavitation- Erosion Resistance of High-chromium Cast Steel- Preliminary Results”, Wear vol. 225-229, 1999, pp 517-522. 4. F. Wantang, Z. Yangzeeng, and J. Tianfu, Y. Mei, “Structural Changes after Cavitation Erosion for a Cr-Mn-N Stainless Steel”, Wear vol. 205, 1997, pp 28-31. 5. M. Grech and J. M. Young, “Influence of Austempering Temperature on the Characteristics of Austempered Ductile Iron Alloyed with Cu and Ni”, AFS vol. 98, 1990, pp345-352. 6. C. K. Lin, P. K. Lai and T. S. Shih, “Influence of Microstructure on the Fatigue Properties of Austempered Ductile Irons-1 High-cycle Fatigue”, Int. J. Fatigue vol. 18, no. 5, 1996, pp 297-307. 7. J. Aranzabal, I. Gutierrez, J. M. Rodriguez-Ibabe, and J. J. Urcola, “Influence of the Amount and Morphology of Retained Austenite on the Mechanical Properties of an Austempered Ductile Iron”, Metallurgical and Material Transactions A vol. 28A, 1997, pp 1143-1156.

0 20 40 60 80 100 120 140 160Ultrasonic Treating Time (mins)

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8. E. A. Neppiras,”Acoustic Cavitation: an Introduction”, Ultrasonic, 1984, 22 (1), pp. 25-28. 9. Y. Tomita and A.Shima, ”Mechanisms of Impulsive Pressure Generation and Damage Pit Formation by Bubble Collapse”, J. Fluid Mechanics, 1986, 169, pp.535-564. 10. Herman Vetters, “Transformation of Austenite into Bainitic Ferrite and Martensite”, Steel Research vol. 67, no. 10, 1996, pp 408-411.

Fig. 3 Variations of micro-hardness in ADI specimens with ultrasonic cumulative treatment time


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MEETINGS

The Ductile Iron Society 2005 Fall Meeting will be held onOctober 25-26, 2005 in Cleveland, Ohio. More information andonline registration.

BUSINESS

Ashland’s Casting Solutions business taps Taiwan distributorTEMC for alliance

DUBLIN, Ohio (USA) – Ashland’s presence in the Asian foundrymarketplace continues to increase as Ashland Casting Solutions, abusiness group of Ashland Specialty Chemical, a division ofAshland Inc. (NYSE: ASH), announces an agreement with TEMC,Metal and Chemical Corp. of Taiwan. The agreement providesTEMC an exclusive right to distribute Ashland Casting Solutionsbinders and EXACTCAST® sleeve technologies to Taiwanesefoundries.

"TEMC is a leading sleeve and coating supplier for the metalcasting industry in Taiwan. The strong relationship of TEMC withTaiwanese foundries will enable Ashland Casting Solutions to sellproducts in these areas more effectively. We are very pleased tohave TEMC as our exclusive binders and EXACTCAST sleevesdistributor in Taiwan," said Ken Chang, business developmentmanager, Taiwan, Ashland Casting Solutions.

“To be positioned in the Taiwanese market with TEMC is verypositive for both companies,” said Mike Swartzlander, vicepresident, Ashland Specialty Chemical and general manager,Ashland Casting Solutions. “Their reputation in fiber sleeves andcoatings coupled with our leading global position in binders andEXACTCAST sleeves provide a strong offering to foundries in thismarket.”

"We are very excited to have the alliance with Ashland CastingSolutions. Taiwanese foundry operators have known the Ashlandname for more than 25 years, and Ashland means foundrytechnologies, foundry solutions and quality. We are committed toputting 100 percent of our efforts into growing the foundrybusiness with Ashland in Taiwan," said F.K. Chang, generalmanager of TEMC.

Ashland Casting Solutions, a business group of Ashland SpecialtyChemical, is a leader in supplying products, processes andtechnologies to the global metal casting marketplace. The group hasoperations (including licensees and joint ventures) in 21 countries.

Ashland Specialty Chemical, a division of Ashland Inc., is aleading, worldwide supplier of specialty chemicals servingindustries including adhesives, automotive, composites, metalcasting, merchant marine, paint, paper, plastics, watercraft andwater treatment. Visit www.ashspec.com to learn more about these


susan

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operations.

Ashland Inc. (NYSE: ASH) is a Fortune 500 chemical andtransportation construction company providing products, servicesand customer solutions throughout the world. To learn more aboutAshland, visit www.ashland.com.

® Registered Trademark, Ashland Inc.

Ashland’s Casting Solutions to market SELEE ceramic foamfilters

DUBLIN, Ohio – Ashland’s focus on customer solutions is beingfurther enhanced by a marketing agreement between AshlandCasting Solutions, a business group of Ashland SpecialtyChemical, a division of Ashland Inc. (NYSE: ASH), and SELEECorporation. The agreement provides Ashland with exclusive rightsto market SELEE’s ceramic foam filters for foundry applications inEurope and Asia and non-exclusive rights to distribute the productsin North America.

“SELEE’s ceramic foam filters are a great match to our customer-solutions approach to enhancing foundry production. Through thisagreement, we add additional design services capabilities so we cannow more readily work with our customers from concept tomanufacturing,” said Mike Swartzlander, vice president, AshlandSpecialty Chemical, and general manager, Ashland CastingSolutions.

SELEE is a recognized high-quality producer of ceramic foamfilters in the Americas iron and steel foundry market. Theagreement with Ashland will expand that presence within NorthAmerica and globally. Additionally, Ashland’s technical supportservices for the castings industry will expand to cover SELEE’sceramic foam filters.

“This agreement with Ashland strengthens our existing sales andservice network and will help to further establish SELEE productsin the European and Asian regions. The real winners will be thefoundry industry who will now have greater access to the world’shighest quality ceramic foam filters,” said Mark Morse, president ofSELEE Corporation.

SELEE Corporation is a wholly owned subsidiary of Porvair plc.Porvair plc is a specialist filtration and advanced materials businesswith expertise in a wide range of markets including the filtration ofmolten metals. It invented ceramic foam filtration and is a marketleader in the filtration of molten aluminum, iron and steel. Porvairis listed on the London Stock Exchange. More information can befound at www.porvair.com and www.selee.com.

Ashland Casting Solutions, a business group of Ashland SpecialtyChemical, is a leader in supplying products, processes andtechnologies to the global metal casting marketplace. The group hasoperations (including licensees and joint ventures) in 21 countries.

Ashland Specialty Chemical, a division of Ashland Inc., is aleading, worldwide supplier of specialty chemicals servingindustries including adhesives, automotive, composites, metalcasting, merchant marine, paint, paper, plastics, watercraft andwater treatment. Visit www.ashspec.com to learn more about theseoperations.

Ashland Inc. (NYSE: ASH) is a Fortune 500 chemical andtransportation construction company providing products, servicesand customer solutions throughout the world. To learn more about

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