Cleavage fracture of microalloyed forging steels

ScriptaMetallurgicaef Materialia,Vol. 32, No. 3, pp. 395-400,1995 Copyright0 1994Hsevier ScienceLtd Printed in the USA. All rights reserved 0956-716x/95$9.50+ .OO

CLEAVAGE FRACTURE OF MICROALLOYED

FORGING STEELS

M.A. Linaza, J.L. Romero, J. M. Rodriguez-Ibabe and J.J. Urcola ESII de San Sebastirin and CEIT P” Manuel Lardizabal 13-1520009 San Sebastih, Basque Country, Spain (Received June 21,1994) (Revised September 8,1994) Introduction Microalloying in medium C forging steels together with new processes such as thermomechanical treatments (previously developed for low C steels) or direct quenching from the finishing forging temperature (ausforging) are being used to avoid expensive heat treatments (quenching + tempering) and to obtain required optimum mechanical properties directly after forging. Good strength levels have been obtained without difficulty but for the same strength level, toughness is improved in quenched + tempered steels. In recent years, efforts have been concentrated on improving the toughness by refining the microstructure, and for that, Ti addition is becoming a common method (inhibition of grain coarsening by small TiN precipitates). Several recommendations are made to obtain the maximum yield of fine TIN precipitates, but a certain proportion of TiN precipitates is usually in a coarse form. These precipitates act as cleavage nucleation sites [l-2]. The present work reports the influence of TIN particles on cleavage fracture stress in four microalloyed forging steels with ferrite-pearlite microstructures and one with lath martensite structure. The results show that cleavage fracture propagation at liquid nitrogen temperature is controlled by the energy required for the microcrack to trespass the TIN particle-matrix interface. Experimental Procedure The chemical composition of the four steels studied is given in Table I and they were supplied by GSB (P. Echeverria and AFORASA). Ti-V and high Ti steels were obtained by conventional casting in the form of 2500 kg. square ingots. Low Ti and C-Mn-B steels were obtained in the form of continuous cast 130 mm square billets. The steels were hot rolled to 50 mm square bars. The steels were studied in the as-received condition (microstructures in the asrolled condition consisted of ferrite and pearlite) and, in the case of C-Mn-B steel, a martensitic structure obtained by quenching (1 hour at 950°C followed by water quench) was analysed also. TABLE I - Chemical Composition of the Steels in Wt %. Steel Ti - V low Ti high Ti C-Mn-B

C 0.37 0.35 0.23 0.27

Mn 1.45 1.56 1.72 1.20

Si 0.56 0.33 0.23 0.23

P 0.010 0.004 0.011 co.02

S 0.043 0 007 0.009 co.02

v 0.11 -

Al 0.024 0.027 0.023 0.043

Ti 0.015 0.028 0.044 0.064

B(ppm) 27

N (ppm) 162 89 75 80

The cleavage fracture strength was measured testing notched bend bars. The configuration of the specimen is the same as GrifEths and Owen’s specimen (12.7 mm x 12.7 mm x 75 mm with a centrally located 45” notch 4.23 mm deep and 250 pm notch radius) [3]. Tests were conducted at liquid nitrogen temperature (77K) at a constant crosshead velocity of 2. 10J mm/s.

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Results The addition of Ti results in both fine and coarse TiN precipitates, the former inhibiting austenite grain growth up to high temperatures and the latter impairing fracture toughness by acting as cleavage nucleation sites [ 1, 21. Figure 1 shows histograms of the minimum and maximum dimensions of the coarse precipitates in the four steels. The mechanical properties measured at liquid nitrogen temperature are summarized in Table II together with the constants of the Hollomon equation. TABLE II - Tensile Test Properties at Liquid Nitrogen Temperature.

The cleavage fracture stress values, oF, calculated from the load at fracture and the Griffiths and Owen finite element analysis [3] are reported in Table III. Fractography analysis of all the specimens showed that the fracture occurred by cleavage and in the majority of the specimens it was possible to identify the cleavage origin of the fracture by analysing the river marks (Figure 2). In all the cases there was a principal origin and, except for one case corresponding to the quenched C-Mn-B microstructure in which a Fe carbide was identified as a nucleation origin, this origin was always associated to a TIN particle (Figure 2). Secondary origins were observed too, coming from primary origins and in the majority ofthe cases related to TIN particles. The accurate distance, d, from the root of the notch to the position of cleavage initiation was measured. Entering this distance into the stress distribution with distance from notch root given by Grifliths and Owen [3], the local cleavage fracture stress, oP*, was determined. Minimum and maximum dimensions (amin and amax) of TIN particles responsible for cleavage initiation were measured also. On the other hand, the first cleavage facet formed during the propagation of the microcrack nucleated at the TiN particle was identified and measured (Dmin and Dm,) and by using stereographic analysis the disorientation angle between the facet plane and the notch plane was determined. All these values are reported in Table III. By considering bP* local fracture stress and approximating the microcrack nucleated at a TiN particle to an elliptical geometry, the effective surface energy yprn necessary to propagate the microcrack along particle-matrix interface can be worked out [4] from the equation:

where Q is a correction for the ellipticity given by: Q =

The calculated y,, values are given in Table III. If the disorientation B angle between the perpendicular to the first formed cleavage facet and the direction of the applied tensile stress is considered, a new local fracture cleavage stress can be determined (OF*(p) = oF*cos2@. With this stress value, a new ypm(B) effective surface energy was calculated. Similarly, if the first grain boundary is assumed to be the critical step for cleavage propagation, taking into account the measured dimensions of the first facet (Dmia and D ,,,aJ and considering the equivalent form of equation (I), y,,,,,, energy value can be determined. These data are reported in Table III.

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TABLE III-

Steel

PlUCRJREOFSTEELS

Macroscopic and Local Cleavage Fracture Stresses, Microstructural Parameters and Effective Surface Energy Values Measured in Four-Point Bending Specimens Tested at Liquid Nitrogen Temperature

=F

d

@@a)

(um)

C-Mn-B a-pearlite

2182 l 2193 2135 2089 2011 2044 2172 2335 1962 2105 2207 2108 1798 * 1869 1890 1878 1793 1745 1796 1735 1572

C-Mn-B a’

2552 + 3447 3407

Ti - V

low Ti

high Ti

397

OF* WJa)

hi,

a,

(w)

(m)

%m

Ypm@>

( J/m2)

( J/m2)

Dmin (ctm)

Dmax Ymm ( J/m2) (w)

2.9

0

7.9

2.9 5.0 3.3 5.9 2.5 4.0 2.9 1.2 2.6

22 0 15 - 0 0 0 0 16 0 - 10

10.6 22.2 13.4 19.0 15.6 23.3 8.4 18.5 16.1 7.3 14.1

10.6 16.4 13.4 16.5 15.6 23.3 8.4 18.5 13.7 7.3 13.3

13 24 32 28 26

40 44 32 48 30

138 147 175.5 208 164

19

36

147

13 13 28

38 26 30

120.5 119 177

4.7 3.0 1.9 3.0 2.4 4.3 10.0 3.6

4.8 4.0 2.1 3.4 2.9 5.8 14.3 4.3

5 7.5 9.4 16.5 -16.5 0 23 0

21.8 12.5 9.5 14.8 10.2 20.5 41.8 13.8

21.5 12.1 9.0 12.5 8.6 20.5 30 12.7

22 28 28 21

36 62 36 26

141 144 124 111 109 92 133 69

1.0

1.2 0.48

0 0

16.1 5.9

16.1 5.9

275 135 180 190 200 213 243 175 295 250 250

2162 1741 2002 1946 2008 2054 1952 1900 2023 2106 2065

1.0 3.4 1.7 2.4 2.0 2.6 1.0 2.9 2.7 1.2 1.9

235 80 250 237 280 190 82 175

1817 1556 1827 1791 1626 1653 1526 1546

135 102

3165 3018

I1.42 ++

(*) particle not identifiedat the origin

P (“I

-

(+) intergranularfracture zone at the origin

(++) Fe carbide particle Discussion _~

Fractographic analysis of tested specimens shows that cleavage origins are located close to the position of the peak tensile stress. In Figure 3 macroscopic cleavage fracture stress (OF) is compared with the local value (oF*) determined by fractographic analysis and in 75% of the cases the difference between them is lower that 10%. It is worth emphasising that the cleavage fracture process is controlled by a critical value of applied tensile stress as it has previously been reported for different microstructure types [S]. Sizes of TiN particles responsible for cleavage initiation (Table III) can be compared with the size distribution histograms (Figure 1). In the four steels with ferrite-pearlite microstructures, maximum TiN (a,,) particle sizes originating cleavage are always bigger than 2 urn (except in one specimen). This implies that an important part of the total coarse TiN particle distribution does not take part in the nucleation of cleavage fracture. In the same way, amax is in the majority of the cases (except in two specimens) lower than 6 urn. The range of particle sizes originating cleavage fracture for ferrite-pearlite microstructures in the present work, above 1urn, is in agreement with results reported by Tweed and Knott [4] in C-Mn weld metals tested under the same conditions and presenting similar values of fracture stress. Nevertheless, the experimental evidence that in the majority of the cases big particles (‘6 urn) do not take place in cleavage initiation and the results obtained with quenched C-Mn-B steel must be analysed taking into account the strength of TiN particles. As it was previously observed in these steels in the brittle-ductile transition zone [I], large TiN particles can break at lower stresses than small particles as a consequence of a Weibull volume

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effect, If cracks are nucleated in big particles at stress values too low to promote propagation of the crack across particle/matrix interface, cracks will blunt and they will not progress to cause the final catastrophic cleavage failure [5]. This effect, together with the lower probability of finding big particles in the zone corresponding to the maximum tensile stress ahead of the notch, can explain the low probability of particles larger than 6 urn being cleavage nucleators, It is also worth emphasising that in the C-Mn-B steel with martensitic structure the size of the particles responsible for cleavage initiation is smaller than in the rest of the cases. Taking into account that heat treatment has not modified TiN particle distribution, the different behaviour can be attributed to the different strength of the matrix. Due to the higher strength of the martensitic structure, cracking of large TIN particles takes place at stresses lower than those necessary to the dynamic propagation of the crack through the interface (Figure 4 corresponds to this case). In consequence, only cracks nucleated in small particles at high stress values can surmount the energy barrier of the interface and at those low sizes, not only TIN particles, but carbides too can originate cleavage. In Figure 5 OF*@) local fracture cleavage stress is plotted against @am;,. The majority of the data are between the two straight lines corresponding (in agreement with equation (I)) to effective surface energy values of 8 and 20 J/m2 with a mean value of 14 J/m2. The range of data is in good agreement with results obtained with different steels [5-71. Bowen et al. [5], considering that the critical defect size is equated to the coarsest observed carbide size in the microstructure, obtained values of in,,, in the range 9-14 J/m2. For eutectoid steels with different microstructures Alexander and Berstein [6] determined for cleavage nucleated in inclusions ypm mean values between 5 and 13 J/m2 (some data reached to 20 J/m2). The upper bound of published data corresponds to the value of 25 J/m2 obtained by Gerberich and Kurman [7]. Finally, if the traversing of a high angle boundary in the matrix is considered as the step that controls catastrophic failure [2], Figure 6 can be plotted, in which the dimensions of the first cleavage facet formed from the propagation of the crack nucleated at the TIN particle have been related with OF*. In the figure the slopes of the straight lines correspond to an upper limit (if they have not been able to stop the crack propagation, the energy of the boundary must be lower than the estimated one) of ymm (effective surface energy required to propagate the crack across ferriteferrite interface). The majority of the actual data are in the range of 100-210 J/m2. In general, ypn, is considered lower than ~~~ [S], but there are very few experimentally measured ~~~ values. Hahn [9] reports a value of 56 J/m2 which agrees with the upper bound estimation made in the present work, shown in Figure 6. In any case, in these tests carried out at 77K, a grain boundary has been observed preventing brittle propagation. Therefore it seems, as it has been pointed out by different authors [4, 5, 81 that in the brittle fracture regime crack propagation is controlled by the particle-matrix interface and not by the matrix-matrix interface.

Conclusions I23-

Brittle fracture process is controlled by the maximum tensile stress. TiN is the origin of brittle fracture, although in a case in high strength martensitic steels the origin was a carbide. Coarsest particles are not always the origin of cleavage. In these tests carried out at 77K, crack propagation through the interface particle-matrix seems to control the cleavage process. Effective surface energy values between 8 and 20 J/m2 have been measured. An upper limit of grain boundary (matrix-matrix interface) effective surface energy has been estimated in 100 J/m2.

Acknowledgments This work was carried out in collaboration with AFORASA (Azkoitia) under ECSC contract no 7210-Ma!938. The authors wish to thank the Community for the economical aportation and AFORASA for the collaboration. M.A.L. and J.L. R. acknowledge financial support from the Basque Government and the Spanish Government, respectively

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399

References 1. 2. 3. 4. 5. 6. 7. 8. 9.

Linaza, M.A., Romero, J.L., Rodriguez-Ibabe, J.M. and Urcola, J.J., Scripta Metall. et Mater., 29, 451 (1993). Linaza, M.A., Romero, J.L., Rodriguez-Ibabe, J.M. and Urcola, J.J., Scripta Metall. et Mater., 29, 1217 (1993). GritIiths, J.R. and Owen, D.R.J., J. Mech. Phys. Sol., 19,419 (1971). Tweed, J.H. and Knott, J.F., Acta Metall., 35, 1401 (1987). Bowen, P., Druce, S.G. and Knott, J.F., Acta Metall., 34, 1121 (1986). Alexander, D.J. and Bernstein, I.M., Met. Trans., 2OA, 2321 (1989). Gerberich, W.W. and Kurman, E., Scripta Metall. et Mater., 19, 295 (1985) Lin, T., Evans, A.G. and Ritchie, R.O., Met. Trans., ISA, 641 (1987). Hahn, G.T., Met. Trans., ISA, 947 (1984). 100

I

a q Ti-V 0

low Ti

0

high Ti

w C-Mn-B

Fig. 1. Histograms of a) minimum and b) maximum dimensions of coarse TIN particles.

1 notch Fig. 2.

Fracture sequence in a 4 point bending specimen tested at 77K (Ti-V steel): a) initiation site and b) detail of cracked TIN particle at initiation site.

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FRACTUREOFSTEELS

A

3000

g E

2500

*&

2000

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Ti-V

0

low

n

high

Ti

0

C-Mn-B

(a-p)

0

C-Mn-B

(a’)

Ti

1500

1000 1000

1500

2000 0 F

2500

3000

3500

, MPa

Fig. 3. Comparison of macroscopic cleavage fracture stress (0~) with local vahte (OF*) determined by fratographic analysis. Dotted line corresponds t0 OF* = 0.9 oF. 4000

Fig 4. Cracked TIN particle without promoting cleavage propagation in the matrix. C&In-B steel with lath martensite structure.

4000

20 J/m? 14 J/N?

200

J/m=

8 J/m= 3000

-

3000

m

2 E &

2 2000

-

*ir.

2000

b

*fr

b

0

low

Ti

1000

high C-Mn-B I

0 0

Fig. 5. OF*@) values plotted against the reciprocal square-root of the minimum TiN dimension responsible for cleavage initiation

n

0

500

Ti

I

(a-p) ,

I

,

1000

Fig. 6. OF* values plotted against the reciprocal square-root of the first cleavage facet minimum dimension.