CARBIDE STEELS WITH INCREASED CRACKING RESISTANCE

N. N. Maslennikov, M. G. Latypov, and A. A. Shatsov

UDC 621.726.01:620.17

Carbide steels are promising materials for components operating in conditions of dry friction, corrosive media, and hydroabrasive wear. Such steels are produced by powder metallurgy procedures using chargesformed from powdered carbon, iron, alloying elements (or a master alloy) and carbides, followed by pressing, sintering, and heat treatment. The structure of the materials consists mainly of a metallic base and carbides. A fundamental disadvantage of this class of material is its poor ductility. The excellent mechanical properties and wear resistance of these steels arise from the ductile and strong matrix and hard carbide inclusions derived from various elements (Cr, Si, Fe, Ti, etc.). At an approximately 20% carbide phase concentration and above, the technological properties of the materials and their ductility are significantly impaired. In this context, interest has been expressed in examining steels with a lower carbide phase concentration.

The characteristics of carbide steels used in this country have been considered in [1]. Carbide steel compositions have also been developed in various foreign countries [1-4]. Particularly vigorous investigations in this direction are being undertaken in Germany [5-7]. In the references quoted above, the carbide steels were produced by sintering in a protective atmosphere at 1280°C via a liquid phase which develops at the sintering stage due to an eutectic forming in the F e - P - C system. This process has two disadvantages. The high phosphorus concentration (0.6%) has an unfavorable influence on the steel ductility, whereas the fairly high liquid phase content and the large volume changes during sintering complicate the procedure. Trials intended to eliminate these disadvantages have not been very successful [8, 9], either the steel volume change during sintering was large or hot isostatic pressing was needed. In order to overcome these problems, we suggest the addition to the charge composition of highly active powders and particle size composition optimization. We have examined* carbide steels consisting of the ZhGrl carbon steel (1% C) and titanium carbide (TIC; 10-20%), and also the ZhGr0.5N12 nickel steel (0.5% C; 12% Ni) and 5-20% TiC. An equivalent cast steel (TD) (Fig. la) produced by the Swedish Defibrator company was used as a comparison. This steel is superior in terms of its wear resistance to other materials employed as grinding spheres in disc mills. It contains 1.7% C; 16.5% Cr; 2.2% Ni; 0.7% Mo, and 1.7% Ti. Specimens were produced by mechanically blending the initial powder constituents, pressing in steel molds at 400600 N/ram2 and sintering in a vacuum over the 1330 to 1400°C range for 2-4 h. The strength properties were determined on plane specimens (GOST 1497-78), and the impact strength on unnotched specimens (GOST 9454-78). Cracking resistance tests were performed on rectangular specimens that had previously been subjected to fatigue cracking during three-point bend testing (GOST 25506-85). Abrasive wear tests were carried out on cylindrical specimens that were 15 mm in diameter using *The authors thank V. N. Antsiferova (corresponding member of the Russian Academy of Sciences) for assistance and support during the work). Research Institute for Powder Technology and Coating Problems, Republican Powder Metallurgy EngineeringTechnical Center, Perm'. Translated from Metallovedenie i Termicheskaya Obrabotka Metallov, No. 8, pp. 20-23, August, 1993.

0026-0673/93/3508-0455512,50

© 1994 Plenum Publishing Corporation

455

TABLE 1

- Ben-ZhT7

Carbide steel

?'

'

HRC

g/cm 3 Z h G r l + 10% T i C . Grl + 20%TIC ; Z h G r O - 5 N 1 2 q-- 10% TiC ' LGrO.5N12 + 20% TiC

~?

?i[¸

strength I ~ ---(%) I ~ _ Krc' J/cm 2, N/ram 2 N/ram3/2 a,

7,2 6,5 7,0

,59 60 55

21 9 29

3700 I 500 1800 I 320 3500 I 810

1500 730 1550

6,5

56

11

2270J 360

360

~i¸l~'!~ iiii,i~

Fig. 1. Microstructure of TD wear resistant steel (a) and ZhGr0.5N12 carbide steel + 20% TiC (b): a) x320; b) x500. ,1m, m g

,

30

o

/

/

7

I

2

v, rain

Fig. 2. Abrasion resistance of powdered steels: 1) ZhGrl + 10% TiC; 2) ZhGrl + 20% TiC; 3) ZhGr0.5N12 + 8% TiC; 4) is the cast steel. a p i n - d i s k system. The specimen was slid along a complex path to constantly encounter a flesh abrasive surface while rotating relative to the sliding direction. The loading was 0.091 N/mm2 and the particle size of the silicon abrasive that was wetted with water was 79 #m. Metallographic, x-ray spectrographic, and x-ray diffraction methods were used to examine the structure of the carbide steels. The optimum sintering time was determined on the basis of the concentration coefficient of variation V = wrD/C, where D is the variance and C is the mean concentration [10]. After disregarding those points having an x-ray intensity below that of the background or greater than the standard, and which also did not conform to the logarithmic normal law for concentration distributions (31 points out of 302), the coefficient VNi was calculated to be 0.087 + 0.008 for 4 h at 1380°C. This was close to the V values for materials obtained by traditional casting technology. A steel sintered at 1380°C for 3 h had a coefficient V = 0.43 and was characterized by a poor set of properties. Heat treatment for the ZhGrl + TiC carbide steel included quenching from 780°C and annealing in oil at 100-120°C. Table 1 gives the results for the properties of the carbide. 456

TABLE 2 Nickel powder form and type

Carbide steel ZhGrO.5N12 + 5% TiC

Carbonyl, PNK-0T4

ZhGrO.5N12 + 10% TiC ZhGrO.5N12 + 15% TiC ZhGrO.5N12 + 20% TiC ZhGrO.5N12 + 10% TiC

Electrolytic, PNE-1

ZhGrO.5N12 + 20% TiC

Poros-

HRC ity

42 l 3 55 30 3l

3 7 12

56

3

57

4

Note. When the PNK-0T4 grade nickel powder was employed, the carbide steel was sintered at 1330°C for 4 h. For the PNI~-I grade nickel powder, the sintering temperature was 1360°C for 4 h.

tI,

~m

~m~ mg

!

.i/

#

/-t z

2S--

3

2O

\

2

L~"

¢

I0 2

0,07

0,02

Fig. 3

#pO3

H -~

0

la

ea

30 r~c,% (vol.)

FIG. 4

Fig. 3. Relationship between wear resistance (Am) and hardness (H) for carbide steels: 1) ZhGr0.5N12 + 10% TiC; 2) ZhGr0.5N12 + 20% TiC. Fig. 4. Relationship between the dispersed phase interparticle spacing and its volume proportion.

Fig. 5. Interaction diagram for carbide steel with abrasive particles: 1) steel matrix; 2) titanium carbide; 3) an abrasive particle. It should be noted that increasing the carbide phase content from 10 to 20% has a neglible effect on the steel abrasion resistance (Fig. 2), which is in good agreement with the data given in [11]. In those cases where the components are large or intricately shaped, quenching is difficult. In order to produce such components, a steel containing 9% Ni and 4% Co is recommended [12]. However, we have established that increasing the nickel content to 12% in the ZhGr0.5N12 + 20% TiC carbide steel that contains no cobalt restricts the residual austenite concentration (Ar) to no more than 10 to 15%. A typical microstructure for the nickel-containing carbide steel is given in 457

Fig. lb. Cobalt is omitted from the steel composition due to its high cost. We believe that the low A r concentration in the carbide steels under examination is due to the fact that titanium carbide is nota rigidly stoichiometric compound and that part of the carbon in the steel reacts with the titanium when it is addedl A reduction in the austenite carbon concentration raises the point at which the martensite transformation commences and terminates, and consequently, reduces the A r level. Further investigations into the ZhGr0.5N12 + 20% TiC steel indicate the fundamental value of the procedure for obtaining nickel powder and the carbide phase level. The use of nickel carbonyl powder enables the sintering temperature to be reduced by approximately 30°C as compared with using electrolytic nickel powder, to produce nonporous components, but hinders the production of components with a high carbide phase concentration. Due to its excellent activity, nickel carbonyl dissolved rapidly in the steel. This allows the contacts between the titanium carbide particles and between these and the steel particles to be reinforced. Moreover, as a result of nickel dissolution at lower temperatures, a liquid phase is formed that ensures fusion of carbide particles. As a result, a carbide framework is formed that prevents shrinkage (Table 2). No carbide framework was noted when specimens containing approximately 20% TiC were heated rapidly. In order to clarify the formation mechanism for the carbide framework, ZhGr0.5N12 + 20% TiC steel specimens were sintered at 1100 and 1300°C. Irrespective of the nickel powder grade, the residual porosity of specimens sintered at both temperatures varied by no more than 1%. As the temperature was increased from 1100 to 1300°C, the specimens' porosity fell by approximately 5%, that is from 23 to 24 down to 18 to 19%. Consequently, the formation of a carbide framework that avoids any porosity reduction is due to the production of a liquid phase at the early sintering stage before shrinkage is terminated as a result of rapid nickel carbonyl dissolution. At a low carbide phase concentration (5-10%), the nickel powder type has no influence on the components density. This is due to the fact that the framework that prevents shrinkage is formed only where there is contact between some portion of the carbides and when sufficiently sized clusters are formed from the carbides. A condition for forming an infinite cluster is that the concentration of the parent component exceeds the occurrence threshold. As follows from percolation theory, the lowest occurrence threshold for three-dimensional model specimens is 0.12 [13]. It is evident from Table 2 that the density of steels for which nickel carbonyl was employed falls abruptly when the carbide phase concentration is greater than 10%. Several approaches are available for assessing the relationship between the abrasive wear resistance and mechanical properties of steels [14, 15]. The data given in Fig. 3 indicate the accuracy of the equation suggested by Khrushchev [15], in which the specimen weight reduction is inversely proportional to its hardness, and which varies as a function of changes in porosity and component concentrations. The weak influence due to the carbide phase at a concentration greater than 10% on the steel wear resistance can be accounted for as follows. We can assume that the distribution of the hardening phase particles in the carbide steel is uniform. If we consider that they are located in the center of a cubic lattice with a mean titanium-carbide grain size equal to 6/zm, then the distance between them will vary in accordance with the curved relationship given in Fig. 4. Material wear arising from abrasive particles takes place more quickly in the saturation where the particles implanted in the steel do not encounter resistant from hardening phase particles. Under these conditions, the criterion for the transition of a carbide steel to a highly wear-resistant state is the spacing L between the titanium-carbide particles. The wear resistance for a carbide steel falls abruptly if L is less than the dimension range of the abrasive particles implanted in the carbide steel (Fig. 5). As a result of the interaction between the abrasive and the steel, the implanted particles become detached. On the basis of the detached particles' size, their contact range relative to the steel may be examined. As a first approximation, the contact range is approximately d, where d is the size of the detached particle. In the case under consideration, the most probable size of the detached particles is 4/~m. Hence, the criterion for a transition to a highly wear-resistant state is given by L _< d. In other words, if the abrasive particles slide along a dispersed hard phase and do not penetrate deeply into the steel matrix, then material wear is slight and is only weakly dependent on the volume proportion of the carbide. Below are given the results from assessing the size distribution of detached abrasive particles. d,/zm 0.125 0.25 to 0.50 0.60 to 0.75 458

no, % 0.20 30.0 2.2

0.99 to 1.65 2.5 to 3.3 4.0 6.0

13.6 4.0 46.0 4.0

Symbols: d is the particles' diameter, n is their number. It follows from comparing the data given above and in Fig. 4 that in the experiment in question, the critical value of L = d = 4 tzm corresponds to a 9% TiC weight proportion. It has been experimentally shown (Fig. 2) that the wear resistance of an alloy containing < 8% TiC decreases abruptly. However, even a steel with 8% TiC has an abrasive wear resistance that is several times greater than its cast equivalent (Fig. 2). A large proportion of the detached particles (30%) have a 0.25-0.5 tzm particle size. Thus, a subsequent rise in the wear resistance may be anticipated when L = 0.25-0_.5/~m, and which corresponds to a titanittm-carbide volume proportion of 46-41% or a 35-31% weight proportion. A significant increase in the abrasive wear resistance at this TiC concentration is specifically recorded in [11]. Thus, a combination of excellent mechanical properties and abrasive wear resistance is available from powdered nickel-containing carbide steels with a 10 to 20% TiC content, and which have a carbide particle size of about 6 #m. It is preferable to use electrolytic nickel powder in such a steel. The abrasive wear resistance for the carbide steel increases in these situations where the small spacing between the dispersed phase inclusions prevent abrasive particle penetration into the steel matrix. REFERENCES ,

2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.

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