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PERFORMANCE OF PM COMPONENTSDURING DYNAMIC LOADING

Report No. PR–03–#1

Research Team: Diana LadosDiran Apelian

(508) 831 5535 lados@wpi.edu(508) 831 5992 dapelian@wpi.edu

Focus Group: Ulf EngstromRuss ChernenkoffRyan Sun/Craig McManusRenato Panelli (via net)Fred SemelDavid Au

North American Hoeganaes, Inc.Ford Motor Company ChairBorg WarnerMahle Metal Leve S.A. (Brazil)HoeganaesQMP

INTRODUCTION

The use of metal powder sintered components is gaining market share forapplications where high integrity and dynamic loading are design considerations.In the last ten years, the annual growth of PM usage has exceeded 6%. A typicalmodern car contains between 15 to 35 lbs of powder components and thisamount is expected to double in the next ten years. There are indubitableadvantages to PM processing especially when components with intricate, andcost-effective near net shape are requirements 1-5. The lack of post sinteringtreatments, uniform quality, reduced material scrap and environmental“friendliness” are factors that make PM technology such a powerful tool 2,6.

The need for understanding the effect of material characteristics, as well asprocessing variables on the dynamic properties in PM components is critical,especially as new and demanding applications surface. Engine and transmissionparts such as connecting rods, camshafts, parking gears, etc., are produced byPM on a large scale 2,7. Fatigue studies have been reported by many in the past,some with reference to PM, 2, 8-15; the literature however is not clear on themechanisms that control fatigue properties, and a systematic approach isneeded, and thus the reason as to why this project was selected by the PMRCconsortium members.

Considering that over 80% of PM materials presently used are PM iron and PMsteels, it seems reasonable to channel our attention to ferrous systems.

Iron powders are produced by chemical procedures (sponge iron obtained byreduction) or by atomization. In atomization, a stream of molten metal is broken

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into particles by a high-pressure jet of water or gas. The molten metal can beeither a pre-alloyed mixture or an iron melt with other elements beingsubsequently admixed or “glued-on” to the Fe particles. Subsequently, thepowders are compacted to a green body under high pressure, 400-1100 N/mm2,and the green body is sintered (heat-treated) between 1100 and 1300°C. Duringsintering, solid-state diffusion takes place, affecting bonding and the requisitemechanical and physical properties to be achieved. The sintering step is usuallycarried out in an atmosphere of H2, dissociated ammonia, or in vacuum, toprotect the compact from oxidation and to reduce the existing oxides. Theatmosphere may also determine the amount and chemical composition of theinclusions. Recrystallized regions and higher densities are obtained when doublepressing and double sintering is utilized. Another method is to use “warmcompaction”, when both the powder and the compaction tools are heated toapproximately 150°C. This raises the green density by 5% of the initial value dueto increased deformability of the powders 2. For very high densities (fully densematerials) powder forging and dynamic compaction are used. It can be seen thatthere are two independent sets of variables when it comes to the topic ofdynamic properties of PM components: one set addresses the variables thataffect powder size, shape, and metallurgical characteristics; the other set dealswith processing variables which affect pore shape, size, etc.

To understand the fatigue behavior of PM materials it is necessary to fullyunderstand microstructure evolution during processing. The typicalmicrostructure of a single phase PM material is mainly characterized by thepresence of pores. These pores, may be isolated, interconnected, or a mixture(Figure 116).

Figure 1. Schematic representation of a PM material.

The amount of pores can be determined from:

01P r

r-=

3

where P=porosity, r=apparent density, r0=density of bulk materialAn alternative way to describe pore structure is to consider open and closedporosity. The closed porosity is equivalent to the isolated pore structure, whereasthe occurrence of pore channels intersecting the surface may be defined as openporosity. The dependence between open and closed porosity for sintered PM ironand steels can be seen in Figure 2 17.

Figure 2. Open porosity vs. total porosity in PM iron and steels.

In a given structure there are only isolated pores if the total porosity is below5%. In technical applications however this condition is rarely met (commonlyporosity level is around and below 15%), and PM iron and PM steel componentscontain isolated and non-isolated pores.

Further quantification of the pore structure is required such as distribution, shapeand size of the pores. The shape of a pore may be described by metallographic

parameters, and also by the shape factor 22

Ur4F p= , where r=maximum radius of

the pore and U=circumference of pores 18.

Prior investigations of PM steels have shown that porosity, pore size and shape,influence fatigue properties more severely than the chemical composition of thepowder 19. However, considering that most of commercial PM components areternary and quaternary alloys (i.e. Fe-Cu-Ni-Mo-C), the influence of the alloyingelements and their interactions cannot be ignored. Moreover, the phases formedduring cooling from the sintering temperatures, is critical.

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As stated earlier, a systematic study where the key variables aredecoupled is needed to have a mechanistic knowledge base to predictdynamic properties of PM components, and to use the knowledge baseto design alloys for specific set of properties.

LITERATURE REVIEW

In the next part of our report selected parts of the most critical reviews will bepresented; the “fuzzy” areas, as well as the gaps in our mechanisticunderstanding will be emphasized. A critical literature review was carried in1996/1997 timeframe, which has been a useful resource, and we will be referringto this review in the sections that follow15.

• Cyclic strain-stress response of PM materials

In fatigue, in general, plastic strain is the decisive damaging parameter. As aresult, the behavior of the microstructure under cyclic loading, stress-strainhysteresis loops need to be determined and evaluated over the entire life of thecomponent in order to asses the hardening and/or softening tendency of thematerial. However, plasticity concepts in PM materials is different from the oneused in conventional materials, in the sense that significant yielding may takeplace locally in the sintering necks close to the pores even if macroscopicallyelastic conditions are still fulfilled. Decoupling these two types of plasticcontributions (microscopic matrix deformation and macroscopic crack openingdispacement) is a complex process and not much research has been carried outon this topic. However, there are some interesting results related to thehardening/softening of the PM materials, such as those reported by Klumpp 20 onpure and double sintered iron.

Figure 3. Changes in hysteresis loops of PM iron

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at different loading cycles in load control.

From the change of the hysteresis loops as a function of loading cycle, it wasobserved that the total strain and plastic strain (width of the hysteresis loop)increased with increasing number of cycles and also the shape of the hysteresisloop became more elliptical at higher numbers of cycles indicating probably anadditional elongation of the specimen due to a time dependent creep process(Figure 3). It was also observed that comparing the single and double sinterediron specimens, the pores in double sintered are rounder and surrounded by amore ductile matrix of larger grains.

The early stages of fatigue loading were observed to be associated withmicrocracks at pore sites, some interconnecting cracks between adjacent poreswith small spacing distance, and some slip bands. Further stages are mostlycharacterized by significant slip development and finally the number ofmicrocracks increases rapidly and macrocracks form leading to final failure. Acomparison between sintered iron and fully dense material (Armco iron) can beseen in Figure 4.

Figure 4. Plastic strain amplitude vs. number of loading cycle for a doublesintered material and a fully dense material.

The two materials show similar tendencies of softening in all the crackdevelopment stages except in the later stages before the final crack growth leadsto failure. In this stage the fully dense material cyclic hardens, while the sinterediron shows a slight amount of cyclic softening, possibly due to enlargement ofmicrocracks overriding matrix hardening. In the final stage, both materials showsimilar softening behavior due to macrocrack growth leading to failure.

The stress-strain dependence could be described by the equation below 9:

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( )napa k e=s

where k=strain hardening coefficient and n=strain hardening exponent.

k is dependent on density (it increases with increasing density and for higherdensities it approaches the value of fully dense material) while n seems to bealmost density independent (Figure 520). However, for double sintered material,both k and n were found to be larger than for the case of single sintered iron.

Figure 5. Stress vs. plastic strain for a single sintered iron with differentdensities.

Piotrowski et al. 21 studied the cyclic stress-strain response of Fe-1.2Cu using theso-called “load increasing step test”. It was observed that for higher density thesame plastic strain is achieved at higher stress amplitudes. Endurance limits weredetermined from this study as the stress values where the plastic strainamplitude vs. stress amplitude curves turn from linearity to exponential behavior.

A homogeneous Ni-Mo PM steel (pearlitic structure) and a non-homogeneous Ni-Mo PM steel (pearlitic, ferritic and martensitic phases present) and aphosphorous containing PM steel were studied 22. For higher densities, all thematerials showed increased endurance limit and cyclic yield strength. The twoNi-Mo steels presented cyclic softening at lower strain amplitudes and hardeningat higher strain amplitudes, while the P steel cyclically hardens. The samehardening was seen in a porous Fe-Cu-Ni steel, while for a porous Fe-Cu-C PMsteel work softening was observed. Similar behavior was reported for wroughtsteels, an indication that the matrix microstructure controls the dynamic

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response both in wrought and in PM steels 23. Lindstedt and Karlsson24 studiedtwo types of AISi 316 stainless steel, a fully dense PM material and a “cast andwrought” material. For various total stress amplitudes, the PM material showedinitial strain hardening, followed by softening and final failure. The cast andwrought materials displayed a strain softening at all stress amplitudes despite

the higher numbers of cycles to failure (Figure 6). The different behavior wasreasoned considering the differences in local chemical compositions of the twosteels. The PM material contains more oxygen and the oxides are located at thegrain boundaries so that the dislocations are slowed down and the cross slip isfavored. The cast and wrought material has a higher nitrogen content promotingplanar glide and thus allowing softening to occur.

Figure 6. Mean peak stress vs. cyclic reversals for two austenitic steels.

• S-N data, Low cycle fatigue (LCF) and High cycle fatigue (HCF)

Most S-N curves in the literature have been determined under rotating bendingloading, due to typical service conditions and because of convenience. Morerecently, other loading modes such as compression/tension, reverse planebending, and torsion 11,14 have been reported but the number of S-N datagenerated in these loading conditions is significantly less than for rotatingbending. Therefore, more investigation is required in order to have a completeunderstanding of the fatigue behavior of PM materials.

Strain life is the preferred approach considering that in common applications,most of the components undergo structural constraints. Strain life curves aredivided into two regions, 0

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behavior occurs at high loading cycles. Nftr represents the number of cycleswhere the tangents to the strain life curve at N=0 and N=• intersect (Figure 7).

Figure 7. Typical strain life curve.

LCF - Low Cycle Fatigue

Few LCF investigations of PM materials are noted in the literature, see Karlssonet al. 22,24 (Figure 8).

Figure 8. Strain life curves for different densities of a PM Fe-1.75Ni-0.5Mo.

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It can be observed that with increasing porosity, the differences between strainlife curves decrease due to reduced microstructural influence on fatigue. For aNi-Mo PM steel, Nftr are in the 10-1000 range with higher values for higherdensities. Conventional ductile steels have Nftr in the 103-104 range whilequenched and tempered steels, with lower ductility, have much lower Nftr values25. A direct comparison between PM and wrought stainless steels 24, showed thatPM steels reached shorter lives when lower cyclic flow stresses were applied,which led to the idea that different mechanistic processes were dominant in thetwo materials.

HCF - High Cycle Fatigue

Endurance limits are commonly defined as the stress amplitudes for differentloading cycles, when the cycles range from 105 to more than 109. According tothe results in the literature stress amplitudes at cycles more than 108 may beregarded as true fatigue limits.

HCF PM iron

Sintered iron alloys were reported 14,20,26 to have increased fatigue limit withincreasing density in a different way. Most investigations achieved higherdensities by optimizing different parameters in the manufacturing process(compacting pressure, sintering temperature and time. According to 27,28

compacting pressure seem to be the most effective to increase fatigue strength.To correlate fatigue limit with density, open porosity seems a more appropriateparameter than total porosity. In Figure 9 17,26 we can notice a change in slope ofthe fatigue strength-porosity curve, for PM iron, and this happens at thetransition point from closed to open porosity.

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Figure 9. Fatigue strength vs. density/porosity for a PM iron.

There are three regions in this curve. Region I corresponds to mostly closedporosity, cracking occurs in the specimen interior, and it connects isolated poresby transgranular propagation. Region II corresponds to the transition fromclosed to open porosity and cracks nucleate mainly at the specimen surface atisolated pores and pore clusters. Several broken sintering necks were observed.Region III corresponds to pores that are all connected to each other and thematerial is essentially regarded as biphasic, with a matrix phase and a porephase. This is the case of completely open porosity. Cracks are nucleatedsimultaneously at different sites at the specimen surface and the broken surfaceis smooth, with the aspect of a ductile fast fracture region and consists of brokensintering necks.

The relationship between fatigue limit and open porosity for a PM iron part isshown in Figure 10 and it can be seen that it is not a linear function.

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Figure 10. Fatigue strength vs. open porosity for a PM iron.

The relationship between fatigue limit and porosity was studied using loadbearing cross-sectional area. Different equations have been employed and someof them were applicable to spherical pores only. A simple approach At=A0(1-P)was used in 17, where At=load bearing cross-section area, A0=nominal cross-section area, and P=porosity. In 29, a more advanced equation was usedAt=A0(1-4pr2/b2) that accounts for the pore structure but with the limitation forspherical pores only. A more accurate load bearing cross-section area wascalculated using the SEM micrographs of the fracture surface and adding all theremaining broken neck areas. Investigations performed in 30,31 show that actualload bearing cross-section is the most efficient corrective method. Fatigue limitwas found to increase with increasing pore shape factors 32.

Another study 33 determined the endurance limit of pure iron sintered from twodifferent types of powders: ASC (water atomized) and NC (reduced sponge ironpowder) using push-pull and bending fatigue testing at R=-1. Results obtainedfrom S-N curves are shown in Figure 11.

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Figure 11. Fatigue limit for 107 cycles vs. total porosity for two types of PM iron.

Bending loading results showed lower fatigue limits than those obtained fromaxial loading. Similar results were reported when open porosity was used as acomparison parameter (Figure 1233).

Figure 12. Fatigue limit for 105 loading cycles vs. open porosity.

From these results it was deduced that fatigue limit is insensitive to openporosity for tension-compression tests, and thus for axial testing volumeproperties need to be considered (open porosity being the decisive parameter).In bending however, the surface properties are more important (maximum valueof the surface roughness Rt may be considered as the main parameter).

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HCF PM steel

The presence of various microstructures and chemical compositions makes theunderstanding of the fatigue behavior of PM steels difficult. Fatigue strength of aPM steel 1.75Ni-1.5Cu-0.5Mo-0.6C (microstructure: tempered martensite andsorbite islands) is shown in Figure 12a 26 and of another PM steel 2Cu-0.8C(microstructure: pearlite and ferrite) is presented in Figure 13b 34.

a b

Figure 13a. Fatigue strength vs. porosity for a Fe-1.75Ni-1.5Cu-0.5Mo-0.6Cb. Fatigue strength vs. porosity for a Fe-2Cu-0.8C.

Fatigue limits increase with reduced porosity but more in-depth investigationsshowed that porosity alone can not satisfactorily describe the fatigue response ofsuch complex materials. However, in sintered steels, pores play a moreimportant role than in sintered pure iron, because the ability of yielding is morepronounced in the latter. Most of the studies were concerned with theoptimization of the pore shape factor through the sintering process consideringthat rounder pores provide higher fatigue strength. Lower sintering temperaturesand higher compacting pressures and longer sintering times (but not longer the2hrs) were beneficial to the fatigue limits 26. Similar studies by Piotrowsky et al.27 and Lindqvist 35 and Sonsino et al. 36 for Cu alloyed steels showed that rounderpores obtained by higher sintering temperature (from 1120°C to 1280°C)increased the endurance limit by 10% and cyclic yield stress by 40%. Christianand German 34 analyzed the influence of the initial size of the powder particleson the fatigue strength. They found that for low densities, small particles arepreferable to achieve higher fatigue strength while for higher densities, largerparticles proved to be more beneficial. For low porosity, rounder pores and largepore spacing were favorable for high fatigue limits. This study pointed out thatpore length, pore curvature and pore spacing are the most relevant pore

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features, influencing fatigue endurance limit. Increasing the Ni content from 1.75to 8% on standard Ni, Cu, Mo, and C steels was found 35 to increase the fatiguestrength through matrix strengthening and enclosing pores in the austeniticphase. Initial powder size seemed to have no effect on the fatigue response.Conventional and PM materials use similar alloying techniques. Ni, Cu, Mo, P, andC, proved to be favorable alloying elements. Increasing the amount of P, Cu, Niand C has a strengthening effect.

PM powders are mostly fabricated using pre-alloyed powders. However, powdermixtures have the advantage of precise composition selection, and bettercompactability. Segregation can be mitigated by the use of bonded powders andcoated powders 37,38. However, comparing fatigue strength of two Fe-0.5Mo-1.5Cu-1.75Ni alloys (microstructure: “divorced pearlite”, martensite, Ni-reachferrite), one diffusion alloyed and one binder-treated it can be noted that fatiguebehavior is similar.

Surface treatments (especially in bending where cracks are initiated at nearsurface regions) lead to enhanced fatigue responses as it does in wroughtmaterials.

• Crack initiation and early stages of fatigue crack growth

In homogenous single phase and engineering materials cracks usually initiate atdifferent stress raisers such as slip bands, inclusions, precipitates, notches, etc.In PM materials these stress raisers are over-imposed on the pore structure,characterized by pore density, size, shape, and inter-pore distance. Stressconcentration factors and locations are not known for these materials; thusdescribing crack initiation is a difficult process compared to fully dense materials.

First to use LEFM in trying to understand the influence of singular defects onfatigue life of PM materials were Betz and Track 39. It was understood that thegeometrical position of the defects plays an important role and three types oflocations were determined: the interior of the specimen, near the specimensurface, and at the specimen surface.

The effect of pore structure on fatigue crack initiation in PM steels has beenstudied29, and it was observed that for structures having closed porosity theinitiation period of the crack is a function of the ratio of the inter-pore spacing topore size. The number of cycles for crack initiation decreases with increasingporosity. Three reasons were proposed for this behavior: decreasing of the loadbearing cross-section; pores acting as stress raisers; pores acting as crackprecursors.

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Weiss et al. 40 investigated the influence of different artificial defects (holes) withvarious sizes and shapes, on crack initiation and early stages of crackpropagation. This study showed that a critical hole geometry for crack initiationmay be predicted by both material properties and geometrical features of thehole.

Klumpp 20 studied the Fe-1.5Cu PM alloy system along with pure PM iron. For thepure iron the strain increases rapidly with the increasing number of cycles, givingrise to a crack very early on, after only a few cycles. For the Fe-1.5Cu alloys thestrain response depends on the number of cycles (in a more complex way) andmicrocracks appear after a larger number of cycles. Pores and inclusions on boththe specimen surface and in the interior were found as sites for microcracking41,42. Cracking occurred in regions of high plastic deformations near inclusionsand pores. The crack density was much higher than that seen in wrought alloys.It can be deduced that the stress field around the pores influence the crack pathmore significantly than the various phases present.

Comparing crack density as function of number of loading cycles for a fully densePM and a wrought stainless steel, Lindstedt and Karlsson 24 observed that thedensity for the PM materials is three times lower than the one for the wroughtmaterial. For both materials cracks start nucleating at ~10% of the total lifetimeand at half lifetime they both present crack density saturation. Therefore, it isreasonable to assume that completely different mechanisms govern the cracknucleation and growth in PM and conventional materials. Both materials havepredominantly surface nucleation sites, and for the PM material there areadditional pores and oxides acting as nucleation sites. For PM materials, defectssuch as oxides, secondary phase, and graphite islands are situated at the particleboundaries (not randomly distributed as in wrought materials), and this favorscracks formation.

Kubicki43 tried to relate crack initiation to fracture mechanics. The maximumstress intensity factor determines the site of crack initiation, which correspondsto the maximum pore with the maximum plastic zone.

• Fatigue crack growth (FCG) and fracture mechanicsconsiderations

These type of investigations are based on da/dN vs. DK diagrams. DK is thestress intensity factor and is defined as: aYK ps=D , where a=crack length andY=geometrical correction factor related to the finite specimen dimensions.

Crack closure contributions are eliminated by using an effective stress intensityfactor DKeff= DK- DKop instead of DK. These diagrams can be divided in threeregions: near threshold fatigue crack growth, Paris regime, and near “pseudo”

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fracture toughness crack growth. da/dN vs. DK curves for both cast and wroughtmaterials and PM materials are given in Figure 14 14.

Figure 14. FCG curves for PM and conventional steels.

PM and conventional steels show similar thresholds, the reason possibly beingrelated to interconnected pores, because of their larger radii, crack tip bluntingmay occur. However, at higher DK levels PM materials are inferior toconventional materials.

For both sintered iron and steels, density plays a significant role on DK. Higherdensity means better resistance to FCG and this can be seen in Figure 15 28 for ahomogeneous Fe-1.75Ni-0.5Mo-0.5C steel (microstructure: divorced pearlite)comparing FCG behavior for various density levels.

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Figure 15. FCG for a PM steel (Fe-1.75Ni-0.5Mo-0.5C) at different densities.

Regarding the threshold region, less data is available for pure iron compared tosteels. In Figure 15 17, the variation of DKth and DKtheff with porosity is presented.

Figure 16. DKth and DKtheff vs. porosity for PM iron.

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The curves in Figure 16 may be divided in three regions. Region I ischaracterized by closed porosity, which does not affect the threshold behavior;In Region II both DKth and DKtheff decrease with increasing porosity; Region IIIrepresents specimens with mostly open porosity. DKth and DKtheff are small andnot very sensitive to porosity beyond 20%. Similar investigations have been doneon PM steels 44-47and the results show that fatigue thresholds increase with adecrease in porosity. Effective thresholds behave similarly with the amendmentthat the effect of porosity is reduced. However, an explanation based on closureeffects is needed, and non-existent.

Regarding the Paris regime where a linear relationship between crackpropagation rate and stress intensity factor holds: ( )mKCdN

da D= , it was observed

that for different densities almost parallel lines are obtained, which is equivalentto equal values of the Paris exponent, m. Crack growth in PM materials is oneorder of magnitude faster than in conventional steels and Paris exponent, m isslightly larger.

Regarding the upper region of the FCG curve, near the pseudo fracturetoughness, DKc, Fleck et al. 45 observed for a typical distaloy Fe-1.75Ni-1.5Cu-0.5Mo-0.5C an increase in DKc with increasing density. Similar results are givenin 48 for a PM steel Fe-4Ni-1.5Cu-0.5Mo-0.6C. Compared to conventional steels,PM steels have much lower DKc ranging from 20-50 MPa m1/2 depending mostlyon the porosity level rather than chemical composition.

It can be noticed that from all the above reported data that the use of PM steelsfor highly stressed applications is a very delicate problem and certainly moreinvestigation is required.

• Analysis of the crack path and crack tip

To describe the crack path of PM materials two aspects need to be considered:the crack path connecting neighboring pores and crack branching withsimultaneous nucleation of secondary cracks. Bertilsson et al. 49described thecrack path in a 0.5Mo-1.75Ni-0.5C in terms of F, the fraction of the total cracklength, with F=1 for a pore free material. Moreover, it was observed that theangles q (crack deflection from the main crack growth direction), is ~40°independent on the porosity level (Figure 17). Therefore, the way the pores arelinked is similar for all the cracks at various density levels. Near the crack tip, thecrack orientation is ~45° to the nominal tensile axis, which could be explained byconsidering dominant shear modes at the crack tip.

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Figure 17. Schematic diagram showing crack path.

Crack propagation for some other steels can be described through thecoalescence of microvoids in the sintering necks. These microvoids nucleate at,and grow from, microstructural defects 45. A thorough analysis of the fracturesurface of a PM steel 0.75 Ni-1.5Cu-0.5Mo-0.5C revealed three mechanisms ofnecks failure:

1. Ductile fracture mechanics, when the microvoids coalesce, even when thecrack growth rate is less than the inclusion spacing.

2. Shear failure at sites where the crack front meets unfavorable orientedparticles and local shear stress approaches yield shear stress.

3. Ductile failure when the ligament formed by the remaining necks is unableto take the applied load.

The conditions at the crack tip, for the same PM steel are schematicallyreproduced in Figure 18 45.

Figure 18. Schematic of the crack tip conditions.

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The reverse plastic zone size is larger than the average particle size and so theuse of LEFM is permitted. The crack tip radius is only ~10% of the pore radiuswhich means that the pores blunt the crack tip and retard crack propagation.Consequently the stress concentration effects of the pores are negligible incomparison to the stress concentration at the crack tip. Therefore, the fastercrack growth in PM materials in comparison to conventional fully dense materials,may be attributed to the reduction of the actual load bearing cross-section area.

• The prediction of the critical defect size

Generally, geometrical shape and distribution of defects such as inclusions,foreign particles, microstructural inhomogeneities, and pores cannot becontrolled during the sintering process. Therefore, determining a critical size forthese defects is important.

The use of LEFM is not an appropriate tool for small cracks as Kitagawa andTakahashi stated in their work 50. Several investigations 9,51 defined the fatiguelimit of a defect-free, un-notched specimen to be equal to the stress determinedfrom the DKth of a small non-propagating crack emanating from a defect ofcritical size. As a result, the allowable stress can be determined as a function ofthe size of an existing defect. Such a curve in a log-log scale is presented inFigure 19 and it is called a modified Kitagawa plot.

Figure 19. Modified Kitagawa diagram.

Defects smaller than the critical value, acrit do not cause fatigue failure and forthis case allowable stress in independent of the defect size and equal to thefatigue limit. For defects larger than acrit the allowable stress decreases withincreasing defect size. For a>acrit (Figure 19), the behavior can be described byLEFM, which is the inclined part of the curve.

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For the case of porous materials, the size, shape, and distribution of inclusionsand pores play an additional role in determining the critical defect size. Asmentioned earlier, three defect locations can be distinguished: internal, nearsurface and at the specimen surface. Each of these locations will manifest theireffect differently in the Kitagawa diagram, the most detrimental location being atthe specimen surface as seen in Figure 20. These diagrams permit eitherdetermination of fatigue limit if the critical defect size and the DKtheff are known,or the critical defect size when the fatigue limit and DKtheff are known.

Figure 20. Kitagawa diagram showing the effect of defect location.

For PM iron and steels, the experimentally determined critical defect size issmaller than the predicted one17. This could be explained by considering thecontribution of the stress field around the defect. Finite element analyses hasestablished that the size of a surrounding stress field is approximately twice thedefect size and so it is reasonable to assume that the effective defect size istwice the pore/inclusion size. This assumption was used in 52 for a PM steelcontaining Mo, and good agreement was found.

In 53, the critical defect size for different pore shape factors was determined andit was shown that parts with ellipsoidal pores have lower fatigue limits thanspecimens containing spherical pores.

For macrocracks nucleating at or near the surface, the rugosity, Rt is a measureof the largest pore from which cracks start. Therefore, the critical defect size was

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assumed 33 to be directly proportional to Rt: acrit=kRt. However this is consideredto be an approximation because experimental observations show much smallercritical sizes than predicted by this equation (acrit=300mm for high porosity andacrit=100mm for low porosity levels 54 versus ~1mm predicted by this equation).

• Effects of loading conditions

The influence of mean stresses and loading mode is significant for mostengineering applications. The importance of mean stress contribution led to thedevelopment of a convenient way to relate stress amplitude to mean stress, viathe Smith diagram, which is built from the results of different S-N curves.

Smith diagrams for cyclic tension and bending for the most common PM steels,with different chemical compositions and surface heat treatments, werepresented in Beiss 55. Curves are shifted to higher stress amplitudes withincreasing density, in the 6.8-7.1 g/cm3 range (Figures 21 a &b). The slope ofthe upper branches of Smith diagrams, are lowered when heat and surfacetreatments are applied. It can be concluded that an increase in mean stress levelresults in a decrease in allowable stress amplitude for push-pull loading, lowdensity levels, acute notches, and surface regions after heat treatment 55,56.

a b

Figure 21a. Smith diagram for cycling tension for PM iron with 0.45%Pb. Smith diagram for cycling bending for PM iron with 0.45%P.

The effect of sintering process on Smith diagrams was investigated and thesintering temperature appeared to be the most important parameter 55. Moreintensive sintering processes, higher temperature, longer times beneficially affectfatigue behavior.

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The influence of alloying elements on the fatigue behavior, especially fordifferent loading modes has not been sufficiently investigated. From the fewstudies 55,57 that are in the literature, it can be noticed that Cu and Ni (together)have a positive influence on the fatigue strength, Mo is beneficial only if C is alsopresent. Systematic causal relationships are non-existent.

Notches have a negative impact on fatigue properties of all materials. For porousmaterials, a distinction between internal notches (pores) and external notches,need to be made. Generally, the notch effect is described by Kt, the elastic stressconcentration factor. For both axial and bending loading increasing Kt results indecreased endurance limits. This decrease is even more evident in wrought castiron than in nodular cast iron and PM steel even at different density levels(certainly, higher density results in improved fatigue properties) 14. Anunexpected finding of this work was that for increasing Kt, PM and wroughtmaterials have similar endurance limits, and for Kt>2 and r=7.1 g/cm3, PMmaterials seem to be superior to wrought materials.

If the porosity level is increased, the fatigue limit is almost unaffected byexternal notches because pores will already act as internal notches 55 (PM Fe-2Cu-2.5Ni with different densities under bending loading).

The machining of the notch surface seems to have an effect on the fatiguebehavior with increasing Kt. Even if milling has a small effect, drilling andpressing show a strong influence on notch sensitivity and specimens behavior inthis case could be compared to the behavior of the bulk material.

There are almost no test frequency effects on the fatigue strength of PM steel.The most used frequency range for PM applications is 5-200 Hz. However, oneamendment should be made, very low frequencies (for example 0.02 Hz)deteriorate significantly the fatigue properties.

Prior to establishing the objectives of this project, few important remarks need tobe made regarding our understanding of the subject matter based on theliterature. As noted above, many investigations have been carried out to studythe influence of different processing parameters, sintering parameters, andalloying techniques, on the fatigue behavior of ferrous PM materials. However,there is almost no information related to the mechanisms of fatigue and fracturefor PM materials. A mechanistic based knowledge base is lacking. Theinteractions of the porous microstructure with the matrix during dynamic loading,and how the presence of pores influences/changes the behavior of the matrixitself, have not been assessed. Fatigue strength data need to be corroboratedwith fatigue mechanisms. The difficulty comes from the fact that fatigue strength

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is a cumulative quantity; different processes of fatigue damage (and themicrostructural constituents that are contributing to them) are important to beidentified, dissociated, and quantified. In conventional materials, microstructuralinhomogeneities/defects act as crack nucleation points, however in PM, porestructure must be taken into consideration. Critical defect size has beenpredicted by other researchers, but again crack initiation needs to be correlatedwith microstructure constituents, and must be done so mechanistically.Moreover, at high loading cycles, the influence of the remaining ligamentmicrostructure on fatigue response has not been studied (most of the studiesconcentrate on the superimposed pores that control damage to failure). Toestablish such a knowledge base, we need to decouple the two majorcontributors, pore structure and matrix microstructure (composition & phases).We view this project to be a systematic investigation to establish a mechanisticunderstanding of fatigue behavior of PM materials.

OBJECTIVE

The objective of the project is to document and investigate/analyze/characterize/understand fatigue failure in PM components as a function of density/porosityand microstructural phases.

EXPERIMENTAL PROCEDURE

With input from Focus Group members, we decided to select a fully pre-alloyedwater atomized low alloy steel powder (Ancorsteel 4600PF) containing Ni, Mo,and Mn, and an admixed alloy, a modified Ancorloy 2, containing Ni, Mo, withsimilar additions of Mn. For chemical consistency reasons, Cu is not added andMn will be added to the admixed alloy - Table 1. The prealloyed material offershigher hardenability, while the admixed alloy will have higher strength andtoughness.

Table 1. Chemical composition of PM alloys selected for this study.

Composition (wt %)Material

C Ni Mo Mn O Graphite

Ancorsteel4600PF 0.01 1.83 0.56 0.15 0.13

Ancorloy 2modified

orAncorsteel

50HP

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While the admixed powders will offer better compactability, segregation might bean issue which we will need to address. We will consider binder-treated powders(using an organic binder that will break down during sintering) or an advancedthermal “glue-on” process where the alloying elements are attached byintermediate temperature diffusion to the individual iron particles.

To understand the fatigue behavior of these materials, it is necessary to considerand analyze their respective microstructures. The typical microstructure of asingle phase PM material is characterized by the presence of pores (isolated,interconnected, or a mixture of both). The amount of pores can be determinedfrom:

01P r

r-=

where: P=porosity, r=apparent density, r0=density of bulk material

An alternative method to describe pore structure is to distinguish the level(s) ofopen and closed porosity, respectively. The closed porosity is equivalent to theisolated pore structure, whereas the occurrence of pore channels intersecting thesurface may be defined as open porosity. We consider that we have isolatedpores if the total porosity is below 5%.

A further description of the pore structure requires distribution, shape and size ofthe pores. The shape of a pore may be described by metallographic parameters,and also by the shape factor

2

2

Ur4F p=

where: r=maximum radius of the pore and U=circumference of pores.

Prior investigations of PM steels have shown that porosity, pore size and shape,influence fatigue properties more severely than chemical composition19 and thuspore structure will be a significant parameter in this study.

The project is divided into three phases:

Phase I:

A thorough, mechanistic understanding of the effects of pore type/amount andtheir interactions on the fatigue response of PM steels will be carried out – i.e.,the various regimes of porosity on fatigue properties… open porosity to isolatedporosity. Subsequently, Different density levels will be selected and for eachdensity level, we will examine and evaluate a controlled mixture of componentshaving isolated and open levels of porosity.

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Phase II:

Microstructural effects will be studied from two points of view:

• One will be to understand at what porosity level (and what ratioclosed/open porosity) microstructure will take over and become the causeof failure; at high porosity levels microstructure effects will not be evidentand the failure will initiate from pores. In order to “allow themicrostructure to be operationally effective” high density (fully-forged)samples are needed. By decreasing the density we will be able to detectthe transition from a matrix/pore controlled system to a mostly porecontrolled system.

• second will be to compare the behavior of two different microstructures(two cooling rates), and optimize the matrix microstructure for fatigueresistance.

Phase III:

A detailed investigation of fatigue response will be carried out from a statefunction or path dependent perspective; a density level will be selected andseveral different processing methods will be followed to attain the selecteddensity level. The effect of processing route (not the density) on gatiguebehavior will be evaluated. This study will provide an advanced understanding ofthe pore size/shape effects on the fatigue behavior.

Detailed Procedures:

Phase I:

First, we will concentrate on the influence of porosity type and amount. Densitywill be varied from 6.5 to 7.86 g/cm3 - four levels as shown in Table 2. The highdensities above 7.6 g/cm3 will be obtained by hot forming at pressures of 415-1100 MPa (30-80 tsi) via powder forging. A 7.86 g/cm3 will essentially representfully-forged conditions with predominantly isolated, closed pores present.

Density[g/cm3]

6.50-6.70 6.90-7.00 7.30-7.40 7.60-7.86

Poretype/amount

High level ofopen porosity

70% openporosity

&30% closed

porosity

70% closedporosity

&30% openporosity

Low level ofclosed porosity

Table 2. Open/closed porosity as a function of density levels.

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In order to select exact density levels conforming with open/closed porosityrequirements, a set of preliminary experiments will be carried out. Samples withdensities from 6.5 to 7.86 will be produced, by adjusting compaction pressures,and the final microstructures will be analyzed to determine the percentages ofopen and closed porosity. Two methods will be used and the results will becorroborated:

1. Deduce the open porosity level by subtracting the closed porosity (viapycnometry) from the total porosity (from geometry/weightconsiderations),

2. Use image analysis to determine the fractions of closed and open porosityon the surfaces of interest.

In this phase sintering temperature and time will be invariants and the desireddensity and pore structures will be achieved by adjusting compaction pressures.

Molding grades particles will be used with an average particle size in the 70-100mm range.

Phase II:

Cooling media and implicitly cooling rates have a critical effect on the finalmicrostructural and the phases present in the resulting component. Afteraustenitizing (2050°F/30 min) the materials will be quenched. Two differentmedia will be used: one that will provide a high cooling rate ~4°F/sec leading toa predominantly martensitic structure, and another a low cooling rate wheremixture of phases will be obtained, martensite-bainite-pearlite-ferrite. Secondaryheat-treatments will be pursued (martensite tempering) to create stable phases,to improve the mechanical properties, and relieve residual stresses (450°F/1 hr inair).

Phase III:

A specific density level (7.4 g/cm3) will be achieved using five different routes:

1. Particles of smaller peak average size: 65-70 mm (lowercompressibility);

2. Particles of larger peak average size: 100-105 mm (highercompressibility);

3. Normal compaction to 7 g/cm3 followed by a temperature/time sinter;4. Double press/Double sinter;5. Surface densification.

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These different processes will not only allow us to investigate/understand if/whythe fatigue strength is path dependent or not, but also it will provide pores ofdifferent sizes/shapes for each case. Therefore, pore morphology and its impacton the fatigue behavior will also be analyzed.

Powder leanliness is a critical variable, and in this study, the inclusion level for allwill be kept as low as possible (commercially) in order to avoid the occurrence ofa new variable. Surface finish on all the samples tested will be similar.

The sintering atmosphere will be 10%H2-90%N2 or dissociated ammonia.

The lubricating additives will be kept constant for all conditions.

Combining all these conditions we will be able to understand the effects of poreamount/type/size/shape as well as the microstructural features and determinethe optimized conditions for fatigue resistance.

Fatigue Tests:

In this work, cylindrical fatigue specimens are selected (Figure 22 is a schematicdiagram giving the dimensions of the cylindrical fatigue specimens). All fatiguespecimens will be machined and polished as spelled out in ASTM standard E466.

127 (5")

20.32 (0.80")

Ø 12.7 (0.50") Ø 7.62

(0.30”)

25.4 RAD (1.00")

Figure 22. Dog-bone specimens used for the fatigue experiments.

The cylindrical fatigue specimens will be tested in pull-pull sinusoidal loading in aservo-hydraulic Instron 8511 machine (Figure 23). The tests will be carried outunder load control at a unique frequency of 50-60 Hz since the test frequencyhas low influence on the fatigue properties (discussion with focus group). Testingwill be conducted at room temperature and controlled humidity.

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Figure 23. Instron 8511.

For selected conditionss fatigue crack growth work, da/dN vs. DK, on compacttension specimens will be done in accordance with ASTM standard E647.

References:

1. R. German, Powder Metallurgy Science, Metal Powder IndustriesFederation, Princeton, NJ, 1984

2. A. Salak, Ferrous Powder Metallurgy, Cambridge International SciencePublishing, Cambridge, 1995

3. K.S. Narasimhan, Adv. Perf. Mater. 3, 19964. W. B. James, Met. Powder Rep. 46, 19915. W. B. James, Int. J. Powder Matall. 30, 19946. Metals Handbook 9th ed., Vol. 7, Materials Park, OH, ASM, 19847. M. Weber, Powder Metall. Int., 25, (3), 19938. A. Wÿhler, Zeitschrift fÚr Bauwesen, 18609. M. Klesnil and P. Lukas, Fatigue of Metallic Materials, Amsterdam, Elsevier,

198010. S. Suresh, Fatigue of Materials, Cambridge, Cambridge University Press,

199111. R. Haynes, Powder Metall., 13, (26), 197012. B. Kubicki, Sintered machine elements, Hertfordshire, Ellis Horwood, 199513. H.R. Jhansale, Proc. Conf. “International Powder Matallurgy Conference”,

(ed. P.U. Gummeson and D. Gustafson), Orlando, FL, June 1988, MPIF,Princeton, NJ

14. E.J. Esper and C.M. Sonsino, Fatigue design for PM components,Shrewsbury, The European Powder Metallurgy Association, 1994

15. A. Hadrboletz and B. Weiss, Fatigue behavior of iron based sinteredmaterial: a review, International Materials Reviews, Vol. 42, No. 1, 1997

30

16. U. Lindstedt, The deformation behavior of PM steel, thesis, Department ofEngineering Metals, Chalmers University of Technology, Gÿteborg, 1995

17. B. Weiss, R. Stickler, and H. Sychra, Met. Powder Rep., 45, 199018. H.E. Exner, Grundlagen von Sintervorg¨ngen, Berlin-Stuttgart,

Borntraeger, 197819. W. Schatt, Pulvermetallurgie, Sinter- und Verbundwerkstoffe, Heidelberg,

Dr. Alfred HÚthig Verlag, 198620. S. Klumpp, Quasistatisches und zyklisches Verformungsverhalten reiner

und legierter Sintereisenwerkstoffe, thesis, University of Karlsruhe,Karlsruhe, 1992

21. A. Piotrowski and D. Eifler, Proc. Conf. “Advances in powder metallurgyand particulate materials”, (ed. J.M. Capus and R.M. German), SanFrancisco, CA, 1992, MPIF, Princeton, NJ, Vol. 6

22. J. Bertilsson and B. Karlsson, Powder Metall. Int., 14, 199223. C.M. Sonsino, F.J. Esper, and G. Leuze, Powder Metall. Int., 14, 198224. U. Lindstedt and B. Karlsson, Proc. Conf. “Advances in powder metallurgy

and particulate materials”, Paris 1994, MPIF, Princeton, NJ, Vol. 325. B.J. Sandor, Fundamentals of cyclic stress and strain, Madison, WI,

University of Wisconsin Press, 197226. H. Sychra, Einflu¢ der Porosit¨t auf das ErmÚdungsverhalten von

Sinterwerkstoffen auf Eisenbasis, thesis, University of Vienna, Vienna,1989

27. A. Piotrowsky and D. Eifler, Proc. Conf. “Powder metallurgy worldcongress”, Kyoto, July 1993, Japan Society of Powder and PowderMetallurgy

28. J. Bertilsson and B. Karlsson, “Powder metallurgy group meeting”, PowderMetallurgy Group, Buxton, 1986

29. J. Holmes and R.A. Queeney, Powder Metall. Int., 28, (4), 198530. J.Bertilsson and B. Karlsson, Proc. Conf. “Fatigue ‘87”, Charlottesville, VA,

EMAS, 198731. P. Franklin and B.L. Davies, Podwer Metall., 1, 197832. O. Buxhaum and C.M. Sonsino, “Horizons of powder metallurgy PM 86”,

(ed. W.A. Kaysser and W.J. Huppmann), Dusseldorf, Freiburg, VerlagSchmid, 1986

33. L. Ledoux and C. Prioul, Proc. Conf. “Modern development in powdermetallurgy”, (ed. P.U. Gummeson and D. Gustafson), Orlando, FL, 1988,MPIF, Princeton, NJ, Vol. 2

34. K.D. Christian and R.M. German, Int. J. Powder Matall., 31, (1), 199535. B.Lindqvist, Met. Powder Rep., 44, 198936. C.M. Sonsino, G. Schlieper, and W.J. Huppmann, J. Powder Metall. Powder

Technol., 20, (1), 198437. H. Danninger, D. Spoljaric, A. Arakil, and B. Weiss, Adv. Powder Metall.,

13, 1996

31

38. N. Chawla, T.F. Murphy, K.S. Narasimhan, M. Koopman, and K.K. Chawla,Axial fatigue behavior of binder-treated versus diffusion alloyed powdermetallurgy steels, Materials Science and Engineering, A308, 2001

39. W. Betz and W. Track, Powder Metall. Int., 13, (4), 198140. B. Weiss, R. Stickler, and A.F. Blom, “The behavior of short fatigue

cracks”, (ed. K.J. Miller and E.R. de los Rios), Mechanical EngineeringPublications, London, 1992

41. D. Rodzinak and M. Slesar, Powder Metall. Int., 12, (3), 198042. G. Medrikat, Beitrag zur Dauerfestigkeitsuntersuchung von Sintereisen,

thesis, Technical University, Berlin, 197143. B. Kubicki, Proc. Conf. “8 Int. Puvermet. Tagung”, Dresden, 198544. H. Sychra et al., Proc. Conf. “12th International Plansee Seminar”, (ed. H.

Ortner), 1989, Plansee Paper HM 4645. N.A. Fleck and R.A. Smith, Powder Metall., 3, 198146. J. Preitfellner, Das ErmÚdungsverhalten von hochfesten Sinterst¨hlen mit

unterschiedlichen Porosit¨ten, thesis, University of Vienna, Vienna, 199247. I.J. Mellanby and J.R. Moon, Met. Powder Rep., June 198948. K.H. Lindner and C.M. Sonsino, Mat.-wiss. u. Werkstofftechnik, 25, 199449. I. Bertilsson, B. Karlsson, and J. Wasen, “Modern developments in powder

metallurgy”, (ed. E.N. Aquaamd C.J. Whitman), Vol.16., 1985, MPIF,Princeton, NJ

50. H. Kitagawa and S. Takahashi, Proc. Conf. “2nd International Conferenceon Mechanical Behavior of Materials”, ASM, Metals Park, OH, 1976

51. K.J Miller, Mater. Sci. Technol., 9, 199352. D. Spoljaric, H. Danninger, D.L. Chen, B. Weiss, and R. Stickler, Proc.

Conf., “advances in powder metallurgy and particulate materials”, Paris,June 1994, European Powder Metallurgy Association, Vol.2

53. G. Harkegard, Eng. Fract. Mech., 6, 197454. B. Weiss and R. Stickler, Proc. Conf. “PM ‘88”, Orlando, FL, June 1988,

Metal Powder Industries Federation, Vol. 2155. P. Beiss, Z. Werkstofftechn., 17, 198656. C.M. Sonsino, Report FB-170, Fraunhofer Institut fÚr Betriebsfestigkeit,

Darmstadt, 198457. C.M. Sonsino, Report FB-1687, Fraunhofer Institut fÚr Betriebsfestigkeit,

Darmstadt, 1983

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