On the fabrication of steel-wire-reinforced white cast irons

MA?ERIAIS SCIENCE & ENCIMEERIWG

A

Materials Scienceand EngineeringA206 (1996) 104-109

On the fabrication of steel-wire-reinforced white cast irons Ma Qian”, Shoji Haradab, Li Yanxiang”, “Department of Mechanical Engineering, Tsinglnra University, bDepnmnent of Mechicnl Engineering, Kyushu Insriiute

Mao Don&u?’

Beijing 100084, of Technology,

People’s Republic of Chinn Kitaky~rsht 804, Jnpm

Received15 May 1995; in revised form 15 July 1995

Abstract White cast irons reinforced by steelwires exhibit much higher toughnessthan do unreinforcedwhite cast irons. I-Iowever, one of the problems encountered during the fabrication processis the melting of steel wires. On the basis of the observed microstructuresof different kinds of steel-iron interface and the assumptionof diffusion-controlled melting, a model was developedin this work to predict the melting behavior of the steelwiresduring the fabrication process.The derived modelwritten as dR2= d2- 40(1 - qO”)/rot, wasconfirmedby experimentalmeasurements. Thus the selectionof steelwireswith respectto the soaking time of the wire or the wall thicknessof the intended castingscan be approximately estimated.

1. Introduction The poor fracture toughness of white cast irons has been the crucial problem that narrows their applications. Although considerable effort has been made from

both metallurgical and mechanical aspects, the toughness of these alloys remains an unreliable factor for repetitive impact applications. Hence, new strategies that could commercially lead to the significant toughening of these alloys are highly expected. The increasingly developing strengthening theories of composite materials, that of the ductile-phase-reinforced brittle matrix composites in particular [l-4], may have provided an enlightening solution to this problem. In fact, white cast iron is nearly a natural brittle matrix composite material reinforced by ductile phases, becauseeither eutectic

M,C or M,C, is generally continuous from a threedimensional view [5], thus acting as the matrix [6]. Such a natural reinforcement has resulted in a fracture

toughness K,, for the composite of around 21 MPa ml’* in contrast with about 3 MPa m”* for the carbide [7]. From the toughening theory of the brittle matrix composites [l-4] it should be possible that the toughness of these materials could be further improved by introducing continuous ductile fibers or wires into them in

addition

to the discontinuous

in-situ

reinforcements.

Recently, Zhao et al. [8] and Aso et al. [9] have individually confirmed that white cast irons could be toughened to a certain extent using steel wire reinforcements (diameter, q3l-6 mm). The impact value of the reinforced irons is as high as five times that of the base alloys [8,9]. This is undoubtedly a very interesting result, implying broader prospects for white irons considering the very low fabrication cost. However, as preliminary studies, all the experiments conducted in [8,9] dealt only with small specimens (specimens of dimensions 15 mm x 15 mm x 120 mm [8] and specimens of diameter 18 mm and length 90 mm [9]), where the soaking time of the wire in the iron melt was so short that the interface reaction (carbon diffusion and melting) might have been heavily suppressed. In principle, the melting of uncoated steel wires associated with the carbon diffusion from the iron melt into the wire surface could hardly be avoided during the fabrication process. In particular, when the solidification time of the intended castings is sufficiently long, this would become a key problem. So that to go further into this strategy toward the practical design of steel-wire-reinforced white iron castings, the melting behavior of steel wires during the fabrication process needs to be made clear. This also forms the main purpose of the present paper. 0921-5093/96/$15.00 0 1996- ElsevierScience S.A. All rightsreserved SSDI 0921-5093~95~10006-7

Ma

Qian et al. 1 Materials

Science

and Engineering

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104-109

Fig. 1. The interface microstructure between steel wire and white cast iron: (a) diameter of 4.50 mm, 2 min at 1300 “C, as cast; (b) diameter of 8.01 mm, 2 min at 1300 “C, as cast; (c) diameter of 8.01 mm, 5 min at 1300 “C, as cast; (d) diameter of 8.01 mm, 5 min at 1300 “C, as cast, then 900 “C for 2 h, air cooled.

2. Experimental

procedures

2.1. Materids nmz’fabrication process

Commercial medium to low carbon (about 0.23 wt.% C) steel wires of different diameters varying from 2.13 to 8.01 mm are utilized. The nominal composition of the base iron is Fe-3.OC-1.8Cr-1.2Si2.7Mn-O.O4S-0.06P (where the composition here and throughout is given in weight per cent). Each time about 300 g of the iron is put in an alumina crucible (diameter, 30 mm; height, 80 mm), and then melted in a silicon carbide tube furnace at 1300 “C and 1250 “C (+2 “C) respectively. 3 min or so after melting of the iron, the steel wire is directly inserted into the melt and then held for different times (2-15 min) so as to simulate approximately the soaking time of the wire during a practical casting process. Next the crucible is taken out and quickly cooled in air. Specimens for examination are all cut from the middle of the ingot. The retained diameters of the wires can be exactly measured from the etched sections of the specimens owing to the etching difference between the iron and the steel.

To define more accurately the steel-iron interface structure during the fabrication process, austenitic stainless steel (JIS SUS 305 (Japan Industrial Standard); Fe-
3. Results and discussion 3.1. The steel-iron interface microstructure

Fig. 1 depicts the general appearance of the steelwire-white-iron interface area, in which the marked interface layer of the discontinuous white phases has been observed as a common phenomenon with respect

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MCI Qinn et nl. / Materids Science ad Engbzeeritzg A206 (lYY6) 103-109

to different steel wires under the experimental conditions (2-15 min). After being heat treated, the white phases decomposed, and small particles of carbide precipitated (Fig. l(d)). Note also that some martensite needles can be found within them (Figs. l(a) and l(b)); these white phases should apparently be austenite. As the holding time increases, it is found that both the thickness of the discontinuous austenite layer and the fraction of the retained austenite in it decrease owing to the weakened cooling effect of the unmelted steel wire. For steel wires of diameter 8.01 mm, we found that the average thickness of the layer varies from 170 kern (2 min) to 100 pm (5 min), 60 pm (10 min) and 50 pm (15 min). For the same holding time, the smaller the diameter of the wires, the narrower was the thickness of the layer and the lower was the amount of the retained austenite. Noguchi et al. [lo] fabricated steel pipes and grey cast irons. Similarly, a distinct layer of transformed austenite (about 120 pm, with no graphite) can also be readily detected at the interface (see Fig. 3(l) of [lo]), where toward the steel side is the pearlite zone caused by carbon diffusion during the fabrication process. Fig. 2 further demonstrates the steel-iron interface structure, where a uniform layer of austenite also clearly exists at the interface (Fig. 2(a)) between the stainless steel and the ductile cast iron. Moreover, the melted part of the stainless steel, i.e. the area of high chromium iron (II, in Fig. 2(a) or 2(b)), the chromium diffusion area in the ductile iron melt, i.e. the area of the ledeburitic structure following II* (Fig. 2(a)) and the carbon diffusion area in the unmelted stainless steel, i.e. the area of fine carbide particles (IA in Fig. 2(a) or 2(c)) are also well distinguished. Fig. 1 (steel wire-hypoeutectic white iron), Fig. 2 (stainless steel-hypereutectic ductile iron) and Fig. 3(l) in [lo] (steel pipe-hypereutectic gray iron) indicate that the steel-iron interface has a layer of austenite after fabrication. This is attributed to the fact that the unmelted steel component-molten iron interface is an austenite-liquid interface, from which the growth of austenite can always occur so long as the temperature gradient is negative in the. iron melt near the interface area before the supercooling for eutectic solidfication is reached there. The following presents direct proof to support this argument. Fig. 3 shows the stainless steel (JIS SUS 305)-ductile iron interface structure subject to a slow cooling rate after solidification. The ductile iron is also a hypereutectic alloy (4.45% CE; Fe-3.70C-2.22Si-0.30MnO.O26P-O.OlOS). In contrast with Fig. 2(a) no (Cr,Fe),C, was found in Fig. 3, and also the spheroidal graphite appeared only about 120 pm away from the interface. This implies that little melting of the stainless steel occurs during fabrication. As the austenite layer has transformed, Fig. 3 then clearly reveals that the

original austenite layer is composed of columnar grains growing from the unmelted stainless steel at the interface. If the transformed austenite layer in Fig. 3 were a part of the unmelted steel component, it would undoubtedly remain an austenitic structure in the case of JIS SUS 305. In studying the transition of graphite morphologies of cast irons (4.36% CE) using unidirectional solidification technique, one of the present authors and coworkers [l l] found that, when the nodularizing agent was

Fig. 2. The stainless steel-ductile iron interface microstructure: (a) as cast; (b) magnified view of II,; (c) magnified view of I,.

basis of mathematically describing the melting haviour of steel wires during fabrication. 3.2. Forndttion

stainless steel

_ ~9

I -40Wll

Fig. 3. The stainless steel-iron interke subject to a little melting of the steel and a slow cooling rate after solidification.

added to the iron melt during solidification, a pure layer of austenite (about 200 urn, without any graphite) always existed between the area of speherical graphite and that of the vermicular or flake graphite (see Fig. 9 in [ll]). The reason is the same. The eutectic solidification of ductile iron requires larger supercooling than that of either vermicular iron or gray iron. Before such a supercooling is established at the interface area, the austenite can directly grow from the austenite phases that already exist at the interface owing to the negative temperature gradient in the iron melt controlled by the unidirectional solidification condition, The reason why the austenite layer may remain at room temperature can be understood through the cooling rate and the temperature MS at which transformation of austenite to ,martensite starts during cooling. From the recommended MS formula [12], MS for the austenite layer would be generally below zero or around the room temperature with respect to 1.32.1 1 wt.% C (from 1300 “C) and certain alloy elements (Mn and Cr) in this work. On the contrary, the cooling effect occasioned by the unmelted steel component after solidification causes a relatively rapid cooling rate near the interface. Thus the austenite layer or a part of it may remain at room temperature. It has been shown that, the stronger the cooling effect caused by the steel wire, the more austenite is retained owing to the low A+!,. Fig. 4 illustrates the schematic profile of the carbon concentration near the interface during fabrication, in which C,,, C,, and C,, denote the carbon concentrations of the initial iron melt, the iron melt at interface, and the steel wire at interface respectively. The maximum value of C,s equals the solubility of austenite at the holing temperature. Cis 6 C,, < C,. The carbon concentration boundary layer defined in Fig. 4 can help to explain the formation of the austenite layer. Identification of the steel-iron interface structure forms the

be-

of the melting model

Since the melting point of the low carbon steel wires is usually higher than the temperature of the iron melt when it is poured into the mold, it can be generally assumed that the melting of steel wire during the fabrication process is occasioned by the carbon diffusing from the iron melt into the surface layer of the wire, and by means of such diffusion the wire is melted layer by layer until solidification occurs at the interface area. Theoretically, the surface layer of the steel component will melt as soon as its carbon concentration exceeds the carbon solubility in austenite at the holding temperature, being above the eutectic temperature of the cast iron. To simplify the model we consider here the melting of a steel wire that occurs at a fixed temperature in the iron melt of such enormous volume that its solute concentration is constant. Also, as the diffusion of carbon in liquid iron is very rapid, we may approximately assume that complete diffusion of carbon will occur in the liquid iron after a short interval at the holding temperature. In this case C,, z C,, or C IL z constant. Thus the solution to the diffusion equation of carbon in the radial direction of the wire can then be written as [13]

where D is the coefficient of carbon, a is the radius of the wire, C, is the initial carbon concentration in the wire and M,, are the roots of Jo(m,, ) = 0

(2)

in which J,(X) is the Bessel function of the first kind of order zero. Let

A , C

GIL

co

CIS

LE Fig. 4. Schematic illustration of the profile of carbon concentration near the interface.

Alo Qim

et al. ,f Mmerinls

Science

rind Engbleering

A206

(1996)

104-109

Eq. (1) then becomes Jo(P,,r/a) = 1 - 2 ccc w(-P,,‘T) (4) IL 1 II = I P,, J, @‘,I1 where & ,f?,,,II = 1, 2, . . , are the roots of J,(p) = 0. As p,, [14], J,,(/?J/~) and J,(/In) have already been tabulated [15], (C- C,)/(C,, - C,) is then only a function of the two dimensionless parameters T and 1*/n. Note that the melting of the surface layer of the wire occurs as the carbon, concentration in it reaches or exceeds the solubility of the carbon in austenite at the holding temperature (above 1150 “C). To formulate the melting ratio we can theri assume that, as (C- C,)/ (C,, - C,) 3 17~(y~+ 1) is satisfied, the surface layer of the wire within Y/U - 1 (i/a I 1) will be melted after an interval t, where C corresponds to the carbon concentration at TICI, and C,, is the solubility of carbon in austenite at the holding tmeperature. Let T= T, and r/n = q. be the roots of Eq. (4) for (C- C,)/ (C’,, - C,) = qO; thus after t, = T0n2/D the melted surface layer would be (1 - &a. Similarly, after t2 = T,qo2a’/D, the melted part will become (1 - qov, . . >so that we can write the melting time of the wire as

c”--“t

t,,=(l

-t402+q04+.

. .+q,2”‘-‘))~

T,C12

or t, =

(1 - qo2”) T,a’ 1 - qo2 D

Rearranging

(6)

Eq. (6) we finally obtain (7)

where cl, is the retained diameter of the wire, or the effective toughening diameter of the wire; dR = dqo”. d is the initial diameter of the wire or, in the strict sense, the diameter of the wire at which complete diffusion of carbon occurs in the iron melt at interface. T, and q,, are two dimensionless parameters, depending only upon the melting temperature for a given value of qo. D can be estimated from [16] D = 0.47 exp(- 1.6C) exp -

(37000 - 66OOC) RT

(8)

in which the coefficient part was derived by Tibbetts [16] according to the measured diffusivity of carbon in austenite from 750 to 1075 “C up to 1.3 wt.% C by the steady-state method, and the energy part (in calories per mole) was taken from the experimental work of Wells et al. [17] covering the temperature range 7501300 “C up to 1.349 wt.% C. C (wt.%) is the carbon concentration in austenite.

-0

2

4

6

8

10

12

14

16

tm (min) Fig. 5. The experimentally determined d,? vs. I,,, relationships of steel wires: curve A, 1250 “C, diameter of S.01 mm; curve B, 13OO”C, diameter of 8.01 mm; curve C, 1250 “C, diameter of 6.48 mm; curve D, 1300 “C, diameter of 6.48 mm.

3.3 Experitnentnl

uet$cntiotz of the model

Eq. (7) reveals that cl, and t, would follow a parabolic relationship or &’ and t, show a linear relationship. Fig. 5 presents the experimentally determined dR2 vs. t, relationships of the steel wires (diameters, 6.48 mm and 8.01 mm) at 1300 “C and 1250 “C respectively. Within the first 2 min the wires melt rather intensively. However, after that is it seen that c&? vs. t, exhibits a good linear relationship (X = 0.997-0.998) for each case. Also, the slopes of curve A (- 1.7636) and curve C (- 1.6354) are approximately the same at 1250 “C, and so approximately are those for curve A (-1.8111) and curve D (-1.6354) at 1300°C. From Eq. (7) the slope of the linear relationship, i.e. 4D(l - q,‘)/T,, would be the same at a fixed temperature. The measreud difference is probably attributed to the diffusion coefficient D which is treated as a constant for different diameters of the wires or difl’erent experimental conditions. As for the assumption of complete carbon diffusion or nearly constant interface concentration, Fig. 5 has also reflected that this assumption approximately works only in about 2 min under the experimental conditions. However, for a rough estimation of the melting behavior of the steel wires it would be still acceptable to treat ci in Eq. (7) as the initial diameter of the steel wire. On substitution of the slopes - 1.8111 and - 1.6354 in Eq. (7), the melting times of the steel wires of diameters 2.13 mm and 4.50 mm are estimated to be around 2.7 min and 12 min respectively at 1300 “C. This is also in general agreement with our experimental observations. Reviewing our preliminary assumptions that (1) the melting of steel wire in an iron melt with a temprature lower than the melting point of the steel wire is occasioned by carbon diffusion and (2) a constant interface concentration can be approximately expected after a short interval at the holding

hJa Qian et al. / lvlaterials

Science

temperature, we believe that Eq. (7) would be of general use for the melting of a thin steel cylinder in an iron melt. As Fig. 5 has apparently shown that ci, instead of the initial diameter dof the wires should be considered for the design of steel-wire-reinforced white irons, Eq. (7) can then be used to predict the effective toughening diameters of the wires by substituting t, for the soaking time of the wires in the intended castings. It is necessary that measures should be taken to protect the wires from being heavily melted when the solidification time of the intended castings becomes sufficiently long. 4. Conclusions (1) After the fabrication of cast iron with a steel component [steel wire and white iron, steep pipe and gray iron, and stainless steel and ductile iron) the steel-iron interface area is usually characterized by a thin layer of retained austenite or transformed austenite. This layer forms during solidification. (2) The melting behavior of the steel wire in an iron melt of the temperature below the melting point of the wire can be generally described by CtR2= d2 - It,, (I is a constant at a given temperature) during the fabrication process. Acknowledgement Helpful suggestions and comments from the reviewer are appreciated.

and Engineering

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References [I! M.F. Ashby, F.J. Blunt and M. Banister,

Acm

Mefail.,

35

(1989) 1847. [2] B.D. Flinn, M. Ruble and A.G. Evans, Acta Merail., 35 (1989) 3001. [31 M. Bannister, H. Shercliff, B. Bao, F. Zuk and M.F. Ashby, Acta Meiall., 40 (1992) 1995. [41 S. Lane, S.B. Biner and 0. Buck, in K. Upadhya (ed.), High P~rjbrmailc’K Composites, Metallurgical

Metal

ad

Society of AIME,

Ceramic,

Mairis

Warrendale, PA.

p. 13. G.F.L. Powell, hJet. Forum, 3 (1980) 37. PI Meids Handbook, Vol. 1 American Society for Metals, Metals Park, OH, 9th edn., 1978, p. 78. 71 (1992) 103. [71 K.H. Zum Gahr, Z. bletnllkd., Brijiug, PI A.-M. Zhao and Z.-C. Wang, J. Univ. Sb. Trchol. 13 (1991) 123. Sot., [91 S. Aso, M. Tagami and S. Goto, J. Jpn. Foundrymen’s 64 (1992) 123. [lOI T. Noguchi, S. Kamota and M. Sakai, Trans. Am. Foundrymen’s Sot., (1993) 231. [I 11 Y.-X. Li, B.C. Liu and C.R. Loper Trans. Am. Fowxlrymen’s sot., 9s (1990) 483. [I21 K.-E. Thelning, Steel wn iis Heat Treafmeni, Butterworths, London, 1984. p. 100. oj’ D$‘Irsiorz, Oxford University [I31 J. Crack, Tile Mat/xwzatics Press, London, 2nd edn., 1975, p. 69. [I41 A. Gray, G.B. Mathews and T.M. Macrobert, A Treafise oiz Bessel Functions, Macmillan, London, 2nd edn., 1922, p. 300. I P51 E.T. Goodwin and J. Staton, Quart. J. Mech. Appl. Mah., (1948) 220. [I61 G.G. Tibbetts, J. Appl. Phys.. 51 (1980) 4813. AI/LIE, IS8 (1950) [I71 C. Welk, W. Batz and F. Mehl, Traw. 553. [51