Wear 254 (2003) 47–54

Comparison of wear resistant MMC and white cast iron

Hans Berns∗Lehrstuhl Werkstofftechnik, Institute fur Werkstoffe, Ruhr-Universität Bochum, IA 2/152, Bochum D-44780, Germany

Received 3 December 2001; received in revised form 28 May 2002; accepted 29 October 2002

Abstract

In this report, the microstructures of conventional white cast irons (WCI) and new metal matrix composites (MMC) are compared. Incontrast to casting, the hot isostatic pressing (HIP) of powder mixtures offers more freedom to design specific properties like toughnessand wear resistance. Both may be enhanced by a proper MMC composition. Experimental results are so convincing that special industrialapplications appear to be feasible as well as cost-effective and some have been initiated. New MMC for high temperature wear, forcorrosive environments and for cold forging tools are presented. It is shown how the cost of hard particles in MMC may be reduced by insitu transformation of ferroalloy particles.© 2002 Elsevier Science B.V. All rights reserved.

Keywords: White cast iron; Metal matrix composite; MMC; Wear resistance; Fracture toughness

1. Introduction

For decades white cast irons (WCI) have been the workhorses of wear protection in the mining and cement industry,as well as in road construction where abrasion by mineralgrains prevails. They are of hypo- to hypereutectic compo-sition and consist of hard carbide phases (HP) embedded ina hardenable metal matrix (MM). Both constituents developfrom the melt and vary in crystallographic structure andhardness depending on alloying and heat treatment.

Recently, wear resistant metal matrix composites (MMC)have been developed, which allow the selection of hardparticles (HP) like carbides, borides, and nitrides and ametal matrix (MM) independently of each other, and de-sign microstructures of superior properties. In contrast tothe solidification of castings in sand moulds close to phaseequilibrium, the powder metallurgical (PM) production ofMMC may stay away from this condition.

In general, the two constituents of wear resistant mate-rials serve different purposes: the HP are to impede wearby grooving or indenting mineral particles while the MM ismeant to provide sufficient toughness. Both properties de-pend on the amount, size and distribution of HP as well ason the hardness and fracture toughness of both constituentsand the bond between them. Because of the comparativelyhigh thermodynamic stability of the HP, material properties

∗ Tel.: +10-49-234-700-5955; fax:+10-49-234-709-4104.E-mail address: berns@wtech.ruhr-uni-bochum.de (H. Berns).

like hot hardness or corrosion resistance are expected torely mainly on the MM[1].

It is the aim of the present study to compare the mi-crostructure and resulting properties of two groups of ma-terials and to explain why the higher costs of MMC maywell pay off. Most results stem from one laboratory to meetconstant test conditions.

2. Microstructure

2.1. White cast iron

The chemical composition and cooling rate determine themorphology of the as-cast microstructure. About 2–4 wt.%of carbon control the amount of HP, while Cr, V, Nb influ-ence their structure and hardness. In addition, Mo and Ni as-sist the hardenability of the MM. The rate of solidification isgiven by the mould material, the cross-section of the castingand the location within. The faster the cooling, the smallerare the primary and eutectic phases. However, the HP areprecipitated from the melt and are coarse compared to thoseformed in solid state. For simplicity let us call primary andeutectic carbides “coarse” and secondary ones “fine”. Thus,the coarse carbide structure is completed at solidus temper-ature, i.e. around 1150◦C at the latest. It remains about un-changed during further cooling or heat treatment (Table 1).

In hypoeutectic irons a eutectic shell solidifies aroundprimary dendrites of MM. In a metallographic sectionthis morphology resembles a carbide net (Fig. 1a). A

0043-1648/02/$ – see front matter © 2002 Elsevier Science B.V. All rights reserved.PII: S0043-1648(02)00300-9

48 H. Berns / Wear 254 (2003) 47–54

Nomenclature

AP abrasive particles (flint, Al2O3, SiC)CWS cold work tool steel (powder, to form

matrix of MMC) (Table 2)DBTT ductile to brittle transition temperature (◦C)F force (N)FDS forging die steel (powder, to form matrix

of MMC) (Table 2)HIP hot isostatic pressingHP hard particles (carbides, borides, nitrides)HSh high speed, high carbon steel (powder,

to form matrix of MMC) (Table 2)HSS high speed steel (powder, to form matrix

of MMC) (Table 2)HV 0.05 Vickers microhardness (F = 0.49 N)HV 30 Vickers macrohardness (F = 294.3 N)HWS hot work tool steel (powder, to form

matrix of MMC) (Table 3)KIc fracture toughness (MPa m1/2)MM metal matrixMMC metal matrix composites (Tables 2 and 4)MSS martensitic stainless steel (powder, to form

matrix of MMC) (Table 4)NiA nickel alloy (powder, to form

matrix of MMC) (Table 2)PM powder metallurgyRb bending strength (MPa)W−1

ab resistance to abrasion (dimensionlesswear resistance)

W−1sa resistance to sliding abrasion

(dimensionless wear resistance)WCI white cast iron (Table 1)

Table 1Grades of wear resistant white cast iron

Designation Mean composition(wt.%)

Coarse(vol.%)

HP type HV 0.05 Retained austenite(vol.%)

HardnessHV 30

W−1ab × 104a

Flintb Al2O3b

WCI 1 Ni2Mo0.5C3.4 48 M3C 1130 0 500 7.4 3.1

WCI 2 Cr12C2 15 M7C3 1450 25 780 4.5 2.7WCI 3 Cr9Ni5Si2C2.9 32 M7C3 1250 15 740 5.4 3.1WCI 4 Cr8Ni5Si2C3.2 38 M7C3 – 20 710 8.4 3.3WCI 5 Cr20Mo1C2.8 30 M7C3 1500 10 790 10 3.8WCI 6 Cr26Mo1C3.2 33 M7C3 1650 <3 790 17 4.1

WCI 7 Cr20Mo6Si3Ni2C3.4 28c M7C3 1700 25 730 10 3.3

WCI 1: near-eutectic alloy with M3C carbides and a perlitic MM, WCI 2: hypoeutectic steel, and WCI 3–6: near-eutectic alloys with M7C3 carbidesof increasing Cr content, hardened and tempered at≤250◦C, WCI 7: as before, but tempered at 515◦C to enhance secondary hardening for elevatedtemperature service.

a Dimensionless abrasive wear resistance (seeSection 3).b 80 mesh abrasive paper.c Plus M6C eutectic of 900 HV.

near-eutectic composition leads to a more homogeneousdistribution of eutectic carbides with occasional primarydendrites or carbides depending on local seggregation. In ahypereutectic alloy the solidification starts with the growthof primary carbides, which are subsequently surrounded bya eutectic net (Fig. 1b). Due to the higher temperature offormation the primary carbides are larger than the eutecticones. In a sand casting of, e.g. 50 mm thickness, the latterare in a 10�m range, while some of the former tend tosolidify in a needle-like shape with an aspect ratio of 10and more, which in thicker parts may extend to the mmrange. As soon as more than one eutectic appears in highalloy irons the constitution becomes more complicated andthe above terminology is no longer applicable.

Starting from an Fe–C melt a grey solidification is turnedinto a white one by increasing the Mn/Si ratio or by chillcasting, to locally obtain a wear resistant surface, while theremainder offers good machineability. The eutectic then con-sists of orthorhombic Fe3C, which reaches a hardness ofabout 1000 HV 0.05 and tends to grow as an interconnectedskeleton with the metal matrix included therein. In this case,MM is somewhat misleading because the HP resembles thematrix of the eutectic in which the MM is embedded. Inpractical application, the composition is near-eutectic, be-cause primary carbides add to the brittleness of the eutec-tic skeleton and primary dendrites lower the resistance togrooving wear. These “unalloyed” irons usually develop aperlitic MM and a hardness of≈550 HV 30.

Next NiCr irons emerge, which are designed to hardenduring cooling in the mould. Here the carbides pick up someCr, turn into M3C of equal morphology and become harder.Ni and Cr increase the hardenability. Ni reduces theAc1 andthe hardening temperature of the MM and lowers its DBTT.As a result of the martensitic and bainitic transformation thehardness increases to≈700 HV 30. Because Ni supportsgraphite formation, the Cr content has to counteract thistendency and the composition is adjusted to the cooling rate,i.e. to the cross-section.

H. Berns / Wear 254 (2003) 47–54 49

Fig. 1. Microstructure of wear resistant materials (schematically): (a) hypoeutectic, high chromium WCI consisting of metal dendrites (MD) surroundedby a net-like eutectic (E) with HP of type M7C3; (b) as (a) but hypereutectic with primary carbides (P) surrounded by a eutectic E; (c) as (a) but withprimary MC carbides (P) by alloying with Nb, V, (Ti); (d) MMC of crushed HP dispersed in a MM; (e) MMC of HP clusters dispersed in a MM (doubledispersion).

The following development targets the increase of carbidehardness and the improvement of carbide morphology byadding from 8 to 26 wt.% Cr at carbon contents between 2and 3.5 wt.%. The resulting M7C3 carbides stem from hexag-onal Cr7C3 but may actually contain more Fe than Cr whichlowers their hardness to the range of 1250–1650 HV 0.05. Atthe lower end of the Cr/C ratio of the alloys the M7C3 maybe surrounded by an M3C layer. The upper limit of this ratiois given by a decrease in the hardenability of the matrix. Incontrast to a eutectic M3C skeleton the M7C3 eutectic con-sists of individual HP embedded in a continuous MM. Thischange in morphology improves the forgeability of alloys onthe low carbon side and the resistance to crack initiation inservice at all carbon levels. Because of a high Cr–C contentthe as-cast MM may consist of up to 100% austenite, whichis decomposed by annealing before hardening and temper-ing. Raising the hardening temperature increases the hard-ness of martensite, but also the austenite/martensite ratio. Atabout 20% of retained austenite a hardness of 850 HV 30 isreached.

These high chromium white cast irons may be furtherimproved by adding Nb, V, (Ti) to precipitate primary cubicMC carbides of superior hardness (2000–3000 HV 0.05).They are of blocky shape and are therefore used to replaceelongated primary M7C3 in hypereutectic CrMo irons. Inhypoeutectic ones they help to protect the MM dendriteswithin the eutectic net (Fig. 1c). They are also a meansof increasing the Cr content of the MM up to the level ofpassivation, i.e. create corrosion resistant irons. For elevatedtemperature service secondary hardening is a means of agingthe as-quenched martensite in the MM by a precipitationof CrMoV-carbides during tempering at 500–550◦C. The

partial exchange of C by B leads to a higher HP hardnessbut is used preferably in hardfacing welding. A number ofWCI are given inTable 1.

2.2. Metal matrix composites

The MMC of Table 2consist of coarse HP dispersed inan MM (Fig. 1d) and were manufactured by hot isostaticpressing (HIP) to reach a fully dense state. The HP werecrushed to an average size of≈80�m, which is roughly anorder of magnitude above the size of eutectic HP in whitecast irons. Agglomerated HP are excluded for lack of in-ner strength. The HP hardness was chosen above that ofthe hardest minerals, namely flint which is a type of quartz(1200 HV 0.05) and corundum (2060 HV 0.05). The fol-lowing HP were used: Cr3C2 (2250 HV 0.05), CrB2 (2500HV 0.05), eutectoid WC/W2C (2600 HV 0.05) (Table 2).Atomized powders of hardenable bcc steels or an fcc Ni al-loy were selected for the MM of MMC as well as for ref-erence specimens without coarse HP. Pure HSh and CWScontain about 30 vol.% of fine globular carbides in the or-der of 1�m in size. About 15 vol.% of these fine HP aredispersed in HSS. Although they are predominantly of eu-tectic origin they fall into the size range of fine secondaryHP because of the rapid quench during atomization andsubsequent globulization during HIP. The Ni superalloy isprecipitation hardened by�′-Ni3Al for elevated temperatureservice.

To obtain a dispersion of the HP in the MM the two pow-ders have to obey a size ratio which depends on their volumeratio (Fig. 2). Relatively fine HP would tend to surround theMM powder grains like satellites forming a brittle net-like

50 H. Berns / Wear 254 (2003) 47–54

Table 2Materials used for PM steels and wear resistant composites

Materials Properties after HIP and heat treatment

Hardness W−1ab × 104

HV 0.05 HV 30 Flinta Al2O3a

MM powder composition (wt.%)FDS Ni1.7Cr1.1Mo0.5V0.1C0.56 660 2.7 2.3CWS Cr13V4Mo1C2.3 740 4.1 (9.1) 2.9 (4.3)HSS W6Mo5V3Cr4C1.3 880 5.0 (10) 3.1 (3.6)HSh W6Mo7V6Co11Cr4C2.3 680 4.8 2.8NiA NiCr20Al4Si3 350 1.8 (2.3) 1.8 (2.1)

HP powders (≈80�m)Cr3C2 2250CrB2 2500WC/W2C 2600

MMC (30 vol.% HP)FDS + CrB2 650 34 4.7CWS + WC/W2C 1040 391 (1735) 21 (62)HSS + WC/W2C 1030 313 19

830 69 9.8NiA + WC/W2C 570 32 (455) 6.7 (28)NiA + Cr3C2 590 42 4.5

FDS: forging die steel, CWS: cold work tool steel, HSS: high speed steel, HSh: as before but of high carbon content, all hardened and tempered, NiA:alloy solution annealed and aged at 750◦C.

a Abrasive paper of 80 mesh (180�m) and in parentheses 220 mesh (67�m).

structure, which has to be avoided. As the HP and MM arenot in thermodynamic equilibrium the HIP conditions de-cide on the amount of interdiffusion between these partners.It is appropriate to keep the pressure high (100–200 MPa)and the temperature low (≤1100◦C) to achieve a goodbond but not too much HP dissolution. At best thin lay-ers of diluted phases like M7C3 around Cr3C2, MB/M3Baround CrB2 or M6C/M12C around WC/W2C are formed(Fig. 1d).

Fig. 2. Microstructure of MMC derived from mixtures of near-globularHP and MM powders of volume contentf and sized. (I, II) Brittledispersion of MM in HP. (III, IV) Ductile dispersion of HP in MM.

3. Properties

3.1. Toughness

This property is important on a microscopic level in thewear surface (microcracking along the rims of wear grooves)and on a macroscopic one (resistance of wear parts to brit-tle fracture). There are a number of ways to measure thetoughness, i.e. the fracture energy of which the notched im-pact test does not differentiate sufficiently. Slow bending ofsmooth specimens is often used for hard materials. However,in the present ones the fracture starts with cracking of the HPand the mechanical properties depend on the length of thesecracks, that means on the size of the HP. This is schemat-ically shown inFig. 3a. The bending strength decreases as

Fig. 3. Schematic representation of bending strengthRb and fracturetoughnessKIc in dependence of the HP size and spacing at a given HPcontent.

H. Berns / Wear 254 (2003) 47–54 51

Fig. 4. Fracture toughness in dependence of the HP content: WCI with(wt.%) 12–26Cr, 2.4Mo, 1.4–3.9C containing M7C3 carbides, tested inthe as-cast, predominantly austenitic state (400–550 HV 50) and afterhardening from 900◦C, deep freezing and tempering at 200◦C to apredominantly martensitic state (700–900 HV 50)[2], MMC as listed inTable 2.

the particle size is raised. Therefore it seems to be reason-able to measure the fracture toughness, which—at a givenHP content—is improved by a growing HP size correspond-ing to a larger HP spacing (Fig. 3b). At a given size of thestressed zone in front of the crack tip, small HP are boundto fracture leading to microcracks ahead of the main crack,which is thereby destabilized. The larger spacing betweencoarse HP may locally confine the stressed zone to the toughmatrix and increaseKIc. Pre-cracked specimens were frac-tured in three-point-bending and compared to theKIc results

Fig. 5. Fracture toughness in dependence of hardness. The near-eutecticwhite cast iron Cr19Mo2.4C2.9 and the MMC FDS+CrB2 contain about30 vol.% of coarse HP, the PM high speed steel HSh about the sameamount of fine ones. There are about 15 vol.% of HP in the hypoeutecticCr12C2 (coarse) and in the HSS (fine). The specimens were hardenedand tempered to the hardness level given.

Fig. 6. Fracture toughness of HP and AP derived from cracks initiatedby Vickers indentations.

of impact tests with sharply notched (EDM cut) specimensreported by[2]. In Fig. 4, the fracture toughness is plottedover the HP content and inFig. 5 over the macrohardness.

While these macroscopic values were measured in bend-ing, the indentation method by Palmqvist and others wasused to derive the fracture toughness of HP and abrasive par-ticles (APs). The eutectoid structure of WC/W2C gives anexceptional high product of fracture toughness and hardness(Fig. 6).

3.2. Wear resistance

To be most effective, reinforcing HP have to be harder andtougher than the AP and at least as large as the groove widthof abrasion. If a wear system changes from erosion-like con-ditions to severe grooving wear, as in mineral extractionand handling, large HP of superior hardness become im-portant. In our pin-on-plate test the end face of a specimen6 mm ∅ was moved back and forth in parallel trails at aspeed of 4.8 mm/s under a normal pressure of 1.32 MPa overfresh abrasive paper while slowly rotating. The dimension-less wear resistanceW−1

ab = ρLA/�m was derived from the

Fig. 7. Wear resistance against 80 mesh flint in dependence of the HPcontent: four MMC are compared to WCI with M7C3, respectively, M3Ccarbides.

52 H. Berns / Wear 254 (2003) 47–54

densityρ, the total lengthL of the wear path (<200 m), thecontact areaA and the mass loss�m. Results of WCI aregiven in Table 1and those of PM matrices and of MMC inTable 2. WCI and MMC with different contents and typesof HP are compared inFig. 7.

4. Comparison

To weigh up different wear resistant materials one has toconsider their expected performance in service as well asthe feasibility of manufacturing and the cost-effectiveness.

4.1. Performance

The target is a high wear resistance combined with suffi-cient toughness for service under conditions of severe groov-ing wear by natural minerals of which quartz is among thehardest. Therefore 80 mesh flint offers a good guideline andfiner or harder abrasives are not considered any further. Com-paring FDS and HSh the 30 vol.% of fine HP (1�m range)of the latter increaseW−1

ab by a factor of only 1.8 (Table 2).The same amount of coarse HP (80�m) in FDS+ 30 vol.%CrB2 yields, however, a factor of 12.6. The factor betweenFDS and WCI 5 containing 30 vol.% of coarse M7C3 (10�mrange) is 3.7 and thus in between the previous two. The besttwo of the MMC are superior to the best two of the investi-gated WCI by a factor of 30–40 (Fig. 7).

A fine HP size and high hardness as revealed by thePM steels HSS and HSh lead to a poor fracture toughness(Fig. 5), which is not compensated by a superior wear re-sistance (Table 2). For HSS the product ofKIc andW−1

ab is55 (omitting 104). In comparison, this product is 950 for theMMC FDS+ 30 vol.% CrB2 and 225 for the near-eutecticWCI 4. In all, the performance in respect to wear resistanceand toughness is considerably improved in the order of hardsteels→ white cast irons→ MMC.

4.2. Feasibility

Compared to casting the size and shape is more difficultto control during HIP which entails more machining. Ingeneral there is a risk of cracking during manufacturingof hard materials. WCI may crack in the mould or duringhardening. The latter may also happen to an MMC espe-cially if it is combined with an unsuited substrate. However,an even dispersion of coarse HP in an MMC is supposed toimprove the crack resistance over an uneven dispersion orinterconnected array of HP in WCI. While the microstruc-ture of WCI depends on the cross-section, that of MMCdoes not. If wear parts have to be partially machined, thevery hard HP in MMC require appropriate machining con-ditions. Joining the MMC to an easily machinable substratemay ease the problem.

The manufacturing costs of PM/HIP are consider-ably above those of sand casting and require a superior

performance of MMC to make this new technologycost-effective. In some applications a change of worn-outparts causes a shut down of a whole plant and thesein-service costs are reduced by a better performance. A highreliability and availability is a key goal of production lines,e.g. in crushing and compaction of minerals. HIP-claddingof tough steel parts adds to the fracture resistance in ser-vice, reduces the costs by limiting the MMC volume andis prone to raise the performance above the WCI level.Crusher rings are a good example of feasible MMC prod-ucts, if performance, manufacturing and costs are jointlyconsidered. HIP facilities of up to about 1.5 m diameterand a loading weight of about 15 t are available in themarket.

5. MMC for special applications

During the solidification of a WCI casting the HP andMM are formed interdependently, which limits the possibil-ities of controlling their individual composition, size, dis-tribution and properties. An MMC allows more variance ofmicrostructural design to achieve specific properties besidesa good abrasion resistance.

5.1. High temperature wear resistance

Secondary hardening of CrMoWV irons (e.g. WCI 7,Table 1) is used to improve the wear resistance in elevatedtemperature service. However, the better hot hardness of theMM is accompanied by eutectic HP. In contrast a matrixof higher hot hardness (HSS, NiA) may be combined withcoarser, harder and more evenly dispersed HP (WC/W2C,Cr3C2) to form MMC of superior high temperature wearresistance. This was confirmed in ring-on-disc experimentswith loose abrasives of 80�m at temperatures up to 900◦C[3–5]. The contact pressure was set at 0.83 MPa and the ro-tational speed at 28 mm/s over a total length of 50 m to givethe dimensionless resistance to sliding abrasionW−1

sa anal-ogous toW−1

ab . While in air wear and oxidation interact, thelatter is excluded in argon atmosphere. Protective layers ofdesintigrated AP, wear debris and oxides were formed onthe wear surface, which increasedW−1

sa up to a tempera-ture TR at which the recrystallization of the MM loweredthe support of the layers. The MMC with a very hard HSSmatrix outperformed the other materials in Ar up to about600◦C, while the maximum wear resistance of the NiA ma-terials was reached at about 700◦C (Fig. 8). The HP contentis most beneficial at 900◦C. In spite of the high Mo con-tent WCI 7 seems to be quite inferior to the MMC. ThreeNiA materials were also studied in air and—in the temper-ature range of 600–800◦C—show a distinctly lower wearresistance compared to runs in Ar. Above 600◦C, WC/W2Cstarted to deteriorate by oxidation while Cr3C2 served wellup to 900◦C. An assessment of our results leads to the con-clusion that MMC offer a new level of wear resistance for

H. Berns / Wear 254 (2003) 47–54 53

Fig. 8. Resistance to sliding abrasionW−1sa of MMC and WCI at elevated

temperatures against 80�m flint.

parts in crushing, sieving, and compacting of hot sinter anddust in the steel industry.

5.2. Corrosion resistance

The strong chemical bond within the HP enhances theircorrosion resistance. The MM has to be protected by passi-vating elements. Powders of hardenable stainless steels with(wt.%) ≥13 Cr and≥0.35 C are suitable. It is well known,though, that the corrosion resistance of stainless martensiteis considerably improved, if the greater part of C is replacedby N [6]. Therefore a (wt.%) Cr15Mo1C0.04 matrix pow-der was mixed with CrN particles. During HIP the reaction2CrN → Cr2N+N yielded N atoms which were dissolved inthe MM and provided a high hardenability and resistance toaqueous corrosion. The passive current density of the MMCCr15Mo1+10 vol.% CrN in 1N H2SO4 was even below thatof the stainless steel grade Cr17Mo1C0.35, which does notcontain coarse HP and wears much faster[7]. Even morepronounced is the resistance of the MMC to pitting by Cl−ions.

5.3. Fracture toughness

In Section 3, it was shown that an increase of the HP sizesimultaneously raises the fracture toughness and the abrasivewear resistance. The bending strength will decrease, becausethe crack length caused by fracturing HP increases (Fig. 3a).However, tools in metal forming require a high strength.This led to replacing the coarse HP, dispersed in the MM,by coarse particles with a high content of fine HP dispersedin an MM (Fig. 1e). This double dispersion microstructure

Table 3Tool life in bolt making

Material of die insert Hardness (HV 30) Numberof boltsa

Increaseby factor

Dieb Wirec

W 6 Mo5Cr4V2C0.9d 640 157 9,640187 2,800

HWS + 60 vol.% HShe 878 157 78,000 8.1187 17,500 6.3

a 12 mm∅, cold headed from wire by die inserts as given.b Hardened and tempered.c Steel with (wt.%) 0.19C and 0.005B annealed to 157 (HV 30) or

cold drawn to 187 (HV 30).d Conventionally cast and hot worked high speed steel.e HWS: hotwork tool steel Cr5Mo1V1Si1C0.4.

may be viewed as clusters of fine HP providing wear resis-tance, surrounded by MM without HP to enhance the frac-ture toughness. It was realized by mixing 60 vol.% of HShpowder of≈80�m size containing about 30 vol.% fine, dis-persed carbides and 40 vol.% hot work steel powder ((HWS)of 16�m mean size. According toFig. 2, and in reality, adouble dispersion microstucture was obtained after HIP. Dieinserts of 33 mm∅ × 12 mm∅ × 10 mm were machined,heat treated and used to form bolts by cold extrusion of steelwire [8,9]. In Table 3, the life of double dispersion die insertsis compared to that of conventional high speed steel, thehardness of which had to be kept rather low to avoid crack-ing. The improvement is quite substantial. This is becausethe fracture toughness of HWS+ 60 vol.% HSh isKIc =16 MPa m1/2 compared to about 10 for HSh (Fig. 5). Doubledispersion or clustered MMC appear to be a means of rais-ing toughness while a high level of wear resistance is main-tained.

5.4. Low-cost HP

The good properties of WC/W2C (Fig. 6) are partly can-celed out by its high price and density, which make an MMCwith 30 vol.% HP rather expensive. Therefore the coarse HPpowder was replaced by a ferroalloy powder (FeTi) with70 wt.% Ti. Upon adding graphite to the MM+FeTi powdermix an “in situ” transformation of the FeTi particles to TiCtook place during HIP. Within a particle the TiC formed ahard case (2600–2750 HV 0.025) and the iron was enrichedin the core (900–1500 HV 0.05). Results of wear tests car-ried out with FDS+ in situ TiC are presented inTable 4and compared to a reference MMC[10]. The higher hard-ness of in situ TiC compared to CrB2 improvedW−1

ab espe-cially against the harder AP used. Compared to pure MMwithout coarse HP (Table 2) the high content of fine TiCin a commercial MMC actually reducesW−1

ab . The reasonis seen in the embrittling effect of the dense population offine HP in front of coarse grooving AP. The material den-sity increases in the order of FeTi, Cr3C2, FDS, WC/W2C.Taking into account the market price for a lot of 250 kg therelative powder cost amounts to 1, 4, 2.3, 50 in the same

54 H. Berns / Wear 254 (2003) 47–54

Table 4Wear resistance of in situ MMC and reference materials against 80 mesh abrasive paper

MMC HP sizea (�m) Hardnessb HV 30 Wear resistanceW−1ab × 104

Flint Al2O3 SiCc

FDS + 10 vol.% in situ TiC 80 733 14.4 5.54 3.14FDS + 30 vol.% in situ TiC 80 720 29.4 7.69 –FDS + 10 vol.% CrB2

d 80 644 11.4 3.70 2.13FDS + 30 vol.% CrB2

d 80 647 33.7 4.72 2.29MSS + 50 vol.% TiCe 3 788 4.07 1.96 1.69

a Mean size.b Hardened and tempered.c Mean hardness (HV 0.05): flint= 1200, Al2O3 = 2060, SiC= 3000.d Volume content includes (Cr, Fe) B rim of interdiffusion.e MSS: martensitic stainless steel Cr13.5Mo3C0.7.

order. In situ MMC are a promizing way to reduce the HPcost.

5.5. Replacement of WCI

The change from ductile Hadfield manganese steel toWCI required more than a decade and was accompanied bynumerous set-backs, i.e. failures in production and service.A similar period is expected for MMC, which have juststarted to enter the market. Small MMC components con-siderably outperformed WCI and hard facing weldments,e.g. in the comminution of hard ores. Up-sizing will taketime. If successful, a gradual replacement of WCI by MMCis conceivable.

Acknowledgements

The work of K. Al-Rubaie, C. Brockmann, M. Buschka,S. Franco, W. Hänsch, S. Koch, O. Lüsebrink, C.V. Nguyen,W. Theisen, W. Trojahn, T. Ümit, B. Wewers carried out atthe Ruhr University is gratefully acknowledged. This invitedpaper was present at COBEM 2001 (Congr. Brasil. Eng.Mec.), Überlandia (MG), 26–30 November 2001.

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[4] H. Berns, S. Koch, High temperature abrasion of a nickel-base alloyand composite, Wear 225–229 (1999) 154–162.

[5] H. Berns, S. Koch, Influence of abrasive particles on wear mechanismand wear resistance in sliding abrasion tests at elevated temperatures,Wear 233–235 (1999) 424–430.

[6] V.G. Gavriljuk, H. Berns, High Nitrogen Steels, Springer, Berlin,1999.

[7] H. Berns, G. Wang, Stainless martensitic PM-HNS, in: Proceedingsof the Conference on High Nitrogen Steels (HNS) 93, Kiev/Ukraine,1993, pp. 415–419; see also Stahl 4 (1994) 38–40.

[8] H. Berns, C.V. Nguyen, A new microstructure for PM toolingmaterial, Met. Phys. Adv. Tech. 16 (1996) 693–706.

[9] H. Berns, A. Melander, D. Weichert, N. Asnafi, C. Brockmann, A.Groß-Wege, A new material for cold forging tools, Comput. Mat.Sci. 11 (1998) 166–180.

[10] H. Berns, B. Wewers, Development of an abrasion resistant steelcomposite with in situ TiC particles, Wear 251 (2001) 1386–1395.