Powder Metallurgy Progress, Vol.11 (2011), No 1-2 104

FATIGUE FAILURE MECHANISM IN Fe-1.4Cr-0.2Mo-1.4Mn-0.7C

M. Kabátová, E. Dudrová, A.S. Wronski, S.C. Mitchell

Abstract Static mechanical properties of Fe-1.4Cr-0.2Mo-1.4Mn-0.7C appear superior to those of Astaloy CrL-0.8C and accordingly their dynamic properties were investigated. Fatigue limit of 190 MPa was found for a stress amplitude of 200 MPa, R = −1 at 24 Hz. Testing was interrupted at 100 cycles and at further 19 intervals until failure after 354300 cycles. For each arrest, surface replicas were made in the two regions where maximum stress was applied. Microcracks could nucleate below 100 cycles, when their sizes ranged from <5 to ~20 μm. Fractographic examination identified the failure-originating site, which was then associated with the crack system observed on the ‘last’ pre-failure micrograph. Detailed examination of the eventual failure region showed nucleation, at various cycle intervals, of 10 microcracks, their subcritical growths, arrests and coalescences with continuing cycling to form a critical crack 3.12 mm deep. The failure path was strongly influenced by defects in the microstructure, associated with probable inadequate processing control of CrL-master alloy-C powder mix.This can account for the fatigue properties being somewhat inferior to CrL -0.7C. Keywords: fatigue failure, manganese steel, intergranular and interparticle fracture, sintering atmosphere

INTRODUCTION There has been extensive recent interest in Cr-Mo base powders, use of master

alloys and manganese steels. Mn is widely used for alloying in wrought steels due to solution strengthening and hardenability. Its affinity for oxygen, however, has posed problems for exploitation by powder metallurgy. Mitchell and Cias [1] examined the thermodynamics of Fe–Mn–C system and concluded that the only effective, practical way to sinter Mn containing steels to be an active ‘microclimate’ around and within the Fe–Mn–C compacts, especially through the use of semiclosed containers. When Salak and Selecka re-examined the problem [2], especially in the context of Mn vapour activity, they concluded that sintering of Mn steels requires a low purity atmosphere, regardless the thermodynamics of the Mn–O system.

Mn had previously been added as ferro-manganese or as elemental powder, but recently interest in master alloys has returned. In particular an European Union Framework 5 project has shown the industrial potential, including prototype manufacture of densified structural parts, of using powder additions of Fe-Mn-C master alloy to iron and iron-base powders, including CrM and CrL [3-6]. Hardness, static and dynamic mechanical properties have been reported for 0.2 [5] and 0.37C [6,7] alloys. In view of the ductility in static mechanical testing persisting to higher C contents, fatigue experiments were now

Margita Kabátová, Eva Dudrová, Institute of Materials Research SAS, Košice, Slovak Republic Andrew, S. Wronski, Stephen C. Mitchell, University of Bradford, Engineering Materials Research Group, United Kingdom

Powder Metallurgy Progress, Vol.11 (2011), No 1-2 105 carried out with 0.7C composition. To evaluate the necessity of a “cleansing” microclimate, the Mn steel specimens were sintered in a commercial-type atmosphere of 10%H2-90%N2 and a dew point no better than -40ºC. The results were to be compared to the, master alloy-free, Astaloy CrL-0.7C specimens sintered at 1120ºC in higher purity atmospheres [8-13].

EXPERIMENTAL PROCEDURES Höganäs Astaloy CrL (Fe-1.5Cr-0.2Mo) was Turbula mixed for 15 mins with a

gas atomised Fe-35Mn-4.4C master alloy (kindly donated by Prof. F. Castro [14]), UF4 graphite and Kenolube, to produce the composition Fe-1.4Cr-1.4 Mn-0.2 Mo-0.8 C. For some of the specimens the master alloy was further milled in a vibratory mill for 15 mins. Bend 40x5x11 (reducing to 5 at centre) mm3 specimens were pressed at 630 MPa and sintered in a closed container containing alumina + 5% graphite powder getter in a laboratory tube furnace in an atmosphere of 10%H2-90%N2 of dew point > -40ºC for 40 mins at 1180ºC. The green density of 6.9 g.cm-3 was essentially unaltered by the sintering process.

To determine the S-N curve, samples were cyclically deformed in simple plane bending on a Schenck fatigue tester, operating at 24Hz with R = -1, For the tests at amplitude S=200 MPa plastic replicas were deposited [8, 9, 15] on both polished faces of the specimens in the areas of the prospective applied maximum stress before testing and after the test was stopped , in the test machine, at selected number of cycles, N. The testing was first stopped after 100 cycles and subsequently 19 times until failure after 354300 cycles. After interrupting the testing and unloading the specimen, replicas were detached and the critical surface areas photographed. Each replica was carbon sputtered and a new replica deposited in the same region; the procedure was carried out at 100, 3000, 10000, 50000, 70000, 100000, 130800, 150000, 200000, 230800, 250000, 271300, 281300, 300000, 307400, 321100, 326600, 330900 and 354300 cycles. 2-3 replicas of the region where failure was expected to be initiated were made.

Using SEM fractography, the failure-initiating region was identified, always at the surface in these experiments. The same region (orthogonal to the fractograph) on both parts of the broken specimen was found, the fracture line photographed and the failure-originating site was then identified on the failure line and on the last pre-failure replica. This same region was examined carefully on all the previously obtained replicas - and photographed. Each microcrack was identified by a capital letter: A-J. Quantitative data were available only for the microcrack axes, “2c”, on the specimen surface; these dimensions were transferred to the fractograph. Thus, at each N, one dimension of each microcrack was known. Careful examination of the microfractographic features and, knowing that the microcrack shapes tend to be near elliptical, enables sketching of the probable microstructure-determined microcrack fronts. Thus the crack depths, “a” axes, were estimated when cycling was interrupted. Then all microcrack sizes (initial, extended and coalesced), i.e. a and c values, were recorded, together with stress intensity factors, estimated using Irwin’s formula [16] for Ka, throughout the fatigue life.

RESULTS The green density of 6.9 g.cm-3 was essentially unaltered by the sintering process.

Following sintering, the microstructure, comprising predominantly pearlite and also upper and lower bainite, is illustrated in Fig.1a and in detail in Fig.1b, where the fine pearlite and slag are evident.

Powder Metallurgy Progress, Vol.11 (2011), No 1-2 106

a) b)

Fig.1. The microstructure of the Fe-1.4Cr-0.2Mo-1.4Mn-0.7C sintered alloy (a) and in detail in (b) where fine pearlite and slag are evident.

Fig.2. A master alloy particle which had not

melted and reacted with the CrL matrix. Fig.3. Partlydiffusion alloyed microstructure

where the transient liquid had not all diffused into the CrL matrix.

Some special features, including those evident on the fractographs, however, need to be reported at the outset. Like in previously reported investigations [4-7] with this master alloy, complete alloying, as illustrated in Figs. 2 and 3, did not take place. To minimise segregation, the master alloy powder must be bound to the base powder [4, 5]. Chromium oxide must be reduced for wetting to take place [4, 5]. In the poor sintering atmosphere (i.e. relatively high dew point) conditions exist for the carbon loss from the master alloy (moving from hyper to hypoeutectic composition), inadequate wetting of oxidised base powder particles and master alloy reactions with slag already present in the base powder. For carbothermic reactions, only a small vapour pressure of water vapour is required and semiclosed containers proved necessary when experiments were carried out in industrial-type furnace [17]. Thus, in addition to the spheroidal isolated pores with the size to ~5 μm, there were major microstructural defects. Further, presence of upper bainite favours fatigue.

Static tensile tests resulted in these values: 0.2% offset yield stress, Rp0.2 = 498 MPa, ultimate tensile strength, Rm = 548 MPa, plastic strain 0.4%, transverse rupture strength, TRS = 905 MPa and fracture toughness K1C = 25-27 MPa.m1/2 (determined in a 3-point bending test, following ASTM Standard E 399-83). The S-N curve, with a fatigue limit of 190 MPa, is represented as Fig.4.

Powder Metallurgy Progress, Vol.11 (2011), No 1-2 107

Fig.4. S-N curve for the tested PM steel.

Tab.1a. Microcrack sizes and stress intensity factors, Ka, at (a) 100<N<130800 (only nucleation and growth).

Number of cycles, N Microcrack A B F C E D c μm 20 a μm 16 100 Ka MPa m1/2 1.8 c μm 75 a μm 50 3000 Ka MPa m1/2 3.0 c μm 100 a μm 70 10000 Ka MPa m1/2 3.6 c μm 100 75 25 50 a μm 70 50 20 46 50000 Ka MPa m1/2 3.6 3.0 2.0 3.3 c μm 135 100 95 35 50 a μm 100 70 80 30 46 70000 Ka MPa m1/2 4.4 3.6 4.1 2.6 3.3 c μm 135 165 95 35 63 a μm 100 120 80 30 54 100000 Ka MPa m1/2 4.39 4.7 4.1 2.6 3.4 c μm 135 215 50 95 35 63 a μm 100 210 46 80 30 54 130800 Ka MPa m1/2 4.4 7.2 3.3 4.1 2.6 3.4

Powder Metallurgy Progress, Vol.11 (2011), No 1-2 108 Tab.1b. Microcrack sizes and stress intensity factors, Ka, at 150000<N<281300 (some coalescences).

Number of cycles, A

microcrack A + B + F G H C E D

c μm 500 95 35 63 a μm 360 80 30 54 150000 Ka MPa m1/2 8.21 4.1 2.6 3.4 c μm 700 75 95 35 63 a μm 450 70 80 30 54 200000 Ka MPa m1/2 8.8 4.1 4.1 2.6 3.4 c μm 925 95 35 63 a μm 550 80 30 54 230800 Ka MPa m1/2 9.5 4.1 2.6 3.4 c μm 1170 95 35 63 a μm 600 80 30 54 250000 Ka MPa m1/2 9.5 4.1 2.6 3.4 c μm 1375 150 100 35 63 a μm 800 120 96 30 54 271300 Ka MPa m1/2 11.4 4.95 4.8 2.6 3.4 c μm 1560 35 63 a μm 1400 30 54 281300 Ka MPa m1/2 17.9 2.6 3.4

Tab.1c. Microcrack sizes and stress intensity factors, Ka, at 300000<N<354300 (ending with growth of the coalesced crack).

Number of cycles, A

microcrack A + B + F + G + H + C + E I J D

c μm 1930 63 a μm 1600 54 300000 Ka MPa m1/2 18.3 3.4 c μm 1960 200 63 a μm 1920 160 54 307400 Ka MPa m1/2 22.0 5.7 3.4 c μm 2500 165 63 a μm 2400 150 54 321100 Ka MPa m1/2 24.2 5.9 3.4 c μm 3085 63 a μm 2962 54 326600 Ka MPa m1/2 26.9 3.4 c μm 3300 a μm 3120 354300 Ka MPa m1/2 27.3

The fracture line (with microcracks identified) and the general fractograph, with

positions of the crack front at various values of N sketched, are presented in Figs.5 and 6. Predominance of inter-particle and inter-grain failure is to be noticed. Microcracking was detected for some specimens examined after testing for only 100 cycles, e.g. Fig.7a, crack

Powder Metallurgy Progress, Vol.11 (2011), No 1-2 109 A, indicating that analyses should consider microcrack initiation from the first cycle. Initial Ka values were as low as 1.8 MPa.m1/2. The sequence of microcrack initiation, arrest, growth, coalescence and propagation is presented in Fig.7a-h, and the relevant crack sizes and Ka values recorded in Table 1.

Fig.5. The fracture line. Identified are the microcracks whose initiatiation, growth, coalescence and propagation resulted in failure.

Fig.6 Fractograph of the fatigue failed specimen of sintered Fe-1.4Cr-0.2Mo-1.4Mn-0.7C specimen, with positions of the crack front at various values of N indicated.

It is seen that 10 microcracks were nucleated at N values ranging from <100 to >307400 cycles. Microcrack nucleation took place at the surface and examples are shown in Fig.8, in which pore clusters and weak particle necks are evident. The first coalescence: microcracks A+B+F, took place at N < 230800. The major coalescence, A+B+F+G+H+C+E+I, to c = 2.5 μm, took place at N> 307400. The final Stage II crack growth was 330900>N>354300, which produced the Stage III propagating crack of c = 3.3 μm and, assuming semi-eliptical geometry, a = 3.12 μm deep. The relevant Ka, 27 MPa.m1/2, approximates to the independently determined K1C of ~26 MPa.m1/2.

Powder Metallurgy Progress, Vol.11 (2011), No 1-2 110

Microcrack growth, Stage II, and eventual propagation, Stage III, paths, of mixed character, were also strongly influenced by defects in the microstructure, Fig.9. In addition to the transgranular paths through fine pearlite, e.g. Fig.9b, of which the cementite lamellae could deflect /arrest the microcracking, microcracks grew along boundaries weakened by poor wetting and/or oxidation, e.g. Fig.9a. Examples of detailed fractographic observations of Stages II and III are presented as Fig.10.

Fig.7. Development of the fatigue producing crack system, (a) to (h) from a 10 micron

microcrack (a). Please note the same area as (h), but after 354300 cycles, has been presented as Fig.6 – at failure.

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(a ) (b)

Fig.8. Location of microcrack initiation sites (a): B at10000<N<50000 and (b): G at 130800<N<150 000.

(a) (b)

Fig. 9. The fatigue crack path exhibits a mixed character well visible on the original surface of the tested specimen (a) and on metallographic etched cross-section (b). (a) shows a crack

path along the original particle surfaces and (b) partially transgranular cracking through fine pearlite with crack deflection by cementite lamellae.

Fig.10(a). Stage II fatigue fracture surface with transgranular facets corresponding to

failure predominantly in pearlite.

Fig.10(b). Stage III failure showing fracture facets resulting from failure along bainitic

packet surfaces and some transgranular facets corresponding to failure of pearlite .

Powder Metallurgy Progress, Vol.11 (2011), No 1-2 112

DISCUSSION The nucleation, growth and coalescence of fatigue microcracks in the sintered

microstructure with ~ 10% porosity, of CrL mixed with Fe-Mn-C master alloy, demonstrates the crucial importanceof clean and careful processing of manganese steels, especially if the manganese is added to a base powder containing chromium, and thus during sintering the particles are covered with a chromium-rich, rather than iron, oxide [18-22]. Indications are that C content, in view of the Mn addition to the already present Cr and Mo, should be no more than 0.5C, [5, 6, 7] in contrast to CrL-C, for which fatigue properties appear attractive [11-13]. The strength of Fe-1.4Cr-0.2Mo-1.4Mn-0.7C is higher than of the unmodified Crl with the same C content, but this is insufficient indication of industrial applicability.

These observations have clearly demonstrated that, when insufficient attention is paid to processing parameters, additional microstructural defects to pores and clean prior particle surfaces affect the fatigue process. In particular interfaces can be contaminated with oxide phases, weak particle necks can form and master alloy particles remain after sintering well above their diffusion and dissolution temperature range. The fatigue process generally resembled the observations of Chawla et al. [23-26] and our previous observations on CrL-C [8-11].

10 microcracks nucleated at various cycles, starting below 100, grew, and coalesced. The joining processes appeared easier through pores and interface areas, especially those contaminated - and this influences the deterioration in fatigue properties. Conversely, microcrack deflection and arrest were due to high-strength obstacles, specifically cementite lamellae in the fine pearlite, e.g. Fig.9b, acting in a similar way to the Ni-rich martensite regions in Fe-0.5/0.85Mo-1.5Cu-1.75Ni-0.6C reported by Chawla et al. [23-24]. Another possible microcrack deflection/stopping mechanism could be the small grains produced by diffusion induced grain boundary migration (DIGM) around pores left by the diffused master alloy particles [28].

Discussion of our findings in the light of recent theoretical analyses, models and simulations [28-35] of failure preceded by coalescence of interacting microcracks, which form a fatal crack that propagates unstably is inappropriate, as clearly the microcrack growth process was step-wise.Accordingly even Paris type analysis is not relevant to Stage II fatigue growth in the PM steels we have investigated [8-11].

CONCLUSIONS 1. These experiments have again clearly demonstrated that, as in static loading of this PM

steel, fatigue failure is the result of growth and coalescence of numerous surface [micro]cracks. Freshly nucleated microcrack, detected within the first 100 cycles, was ~ 20 microns in size. Nucleation, growth and coalescence of 10 neighbouring microcracks were studied in detail by 19 sets of replica microscopic observations. Final coalescence took place to a crack depth of 3.30 mm and the estimate of the then applied stress intensity factor, ~27 MPa.m-1/2, agrees with the fracture toughness value, independently measured as ~26 MPa.m-1/2.

2. As the fatigue (in contrast to static loading) results are inferior to those for CrL+0.7C, it is essential to consider the causes. These must be associated with composition and processing of our specimens, since previous data on CrL+ Fe-Mn-C master alloy +C appear promising.

3. The C content, higher than in the previous investigations, appears too high, 0.5C max is suggested.

Powder Metallurgy Progress, Vol.11 (2011), No 1-2 113 4. The master alloy particles need to be small and bound to the larger base material

powder. Classification to < 20 µm and a binding operation are recommended. 5. To ensure, during the sintering cycle, reduction of the chromium-rich oxide and master

alloy dissolution, careful choice of sintering atmosphere and cycle are essential. 6. Results of dilatometric tests and of static mechanical static properties are insufficient to

fully characterise the processing and properties of manganese containing master alloy modified powders containing chromium.

Acknowledgements The work described forms part of an ongoing collaboration between SAS, Košice

and the University of Bradford. M.Kabátová and E.Dudrová wish also to thank the Slovak National Grant Agency (Project VEGA 2/0103/09) for financial support.

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