2. Materials and Methods
As materials in the CG and UFG states, we used known commercial materials (steels and alloys) that have bcc lattices: certified C45 carbon steel, 9MnSi5 low-alloyed pipe steel, 12Cr-2W-2Ni-0.5Mo ferritic–martensitic steel (
Table 1); hcp lattices: Grade 4 Ti, Ti-6Al-4V titanium alloy, a Mg6Al magnesium alloy (
Table 2); fcc lattices: Fe-0.02C-18Cr-8Ni corrosion-resistant austenitic steel (
Table 1), and an ENAW-2024 aluminum alloy (
Table 3). The chemical composition of the 9MnSi5, 12Cr-2W-2Ni-0.5Mo and Fe-0.02C-18Cr-8Ni steels, as well as of the Ti-6Al-4V and ENAW-2024 alloys, was determined using a Q4 Tasman optical emission spectrometer (Bruker Elemental GmbH, Karlsruhe, Germany). The chemical composition of the Mg6Al magnesium alloy was determined using a Thermo Scientific ARL Optim’X X-ray fluorescence spectrometer (Thermo Fisher Scientific, Ecublens, Switzerland). Comprehensive studies in which the authors produced ultrafine-grained states in these materials via SPD were performed in recent years at the Institute of Physics of Advanced Materials, Ufa State Aviation Technical University (currently Ufa University of Science and Technology, Ufa, Russia) (see [
5,
17]). The specific details of this process are given below.
The C45 steel in the CG state was investigated after quenching and high tempering (550 °C). The UFG structure was produced using equal-channel angular pressing (ECAP) via the following regime: water quenching from a temperature of 800 °C and ECAP at 350 °C (route Bc,
n = 6, φ = 120°) [
5]. The 9MnSi5 steel was studied in the initial hot-rolled CG state and in the UFG state produced using ECAP. The UFG state was produced via the following regime: homogenizing annealing at 820 °C followed by water quenching and tempering at 350 °C, ECAP at 300 °C, and 4 passes via route Bc and additional annealing at 350 °C with a holding time of 10 min. The CG state of the 12Cr-2W-2Ni-0.5Mo steel was produced by heating it in a hot-rolled state to a temperature of 1050 °C for 1 h, followed by oil quenching and tempering at a temperature of 800 °C for 1 h. To produce a UFG state, the steel, after the above-mentioned heat treatment, was processed via ECAP at a temperature of 550 °C (route Bc,
n = 4, φ = 120°). Then, the samples were subjected to additional annealing at a temperature of 850 °C for 1 h plus oil quenching [
18].
Grade 4 Ti in the CG state was investigated after the homogenizing annealing of the hot-rolled billets at a temperature of 680 °C for 1 h. The UFG state of Ti was produced through treatment via the following regime: the homogenizing annealing of the billet at 680 °C for 1 h plus ECAP at a temperature of 250 °C (route Bc, n = 6) and annealing at 600 °C (1 h). The Ti-6Al-4V titanium alloy was studied in the initial CG state, as well as in the UFG state. The alloy’s initial state was produced via hot rolling of the billets. To produce a UFG state, the billet was subjected to homogenizing annealing at 960 °C followed by water quenching and tempering at 675 °C for 4 h and ECAP at 650 °C (route Bc, φ = 120°, n = 6). The Mg6Al magnesium alloy was studied in the initial state and in the ECAP-processed state. The initial state was produced via the homogenizing annealing of the as-cast alloy at a temperature of 430 °C for 24 h in argon medium. After the above-mentioned heat treatment, the Mg6Al alloy was subjected to ECAP at a temperature of 400 °C (n = 4, route Bc, φ = 120°) with intermediate annealing at the same temperature for 15 min.
The Fe-0.02C-18Cr-8Ni austenitic steel was investigated in the initial (hot-rolled) CG state and in the UFG state after ECAP processing. The UFG state was produced after the following treatment: quenching from a temperature of 1050 °C with preliminary holding for 1 h plus ECAP at a temperature of 350 °C (route Bc, n = 4, φ = 120°). The ENAW-2024 aluminum alloy was studied after the standard T6 treatment as follows: heating to a temperature of 530 °C, holding at this temperature for 1 h, and water quenching and aging at a temperature of 190 °C for 7 h with water cooling. The UFG state of the ENAW-2024 alloy was produced after ECAP processing at a temperature of 160 °C (route Bc, n = 6, φ = 90°).
A structural study of the materials with mean grain size evaluation was performed via scanning (JSM-6490LV, JEOL Ltd., Tokyo, Japan) and transmission electron microscopies (JEM-2100, JEOL Ltd., Tokyo, Japan). Hardness tests of the materials were conducted using a TH 300 hardness tester (TIME Group, Beijing, China). The static tension of the cylindrical samples with a diameter of 3 mm was carried out using an H50kT universal testing machine (Tinius Olsen, Redhill, UK) at a temperature of 20 °C, in accordance with GOST 1497-84. Fatigue tests were carried out on prismatic samples 10 mm thick, 15 mm tall and 85 mm in length, with a V-shaped stress raiser and 0.25 mm radius at the vertex. Fatigue tests were performed at a temperature of 20 °C using the three-point bending principle (
Figure 1) using an Instron 8802 system at a loading frequency of ⱱ = 10 Hz, a stress ratio of R = 0.1 and several values of the cycle stress range ΔP. Depending on the investigated material, ΔP varied from 800 to 7000 N. The loading cycle was sinusoidal. Based on the test results, the total number of loading cycles to sample failure (sample life) was found. Upon analyzing the “number of loading cycles–fatigue crack length” curve, the number of loading cycles to fatigue crack initiation was found.
Microfractographic studies of the fracture surfaces were conducted using a Sigma scanning electron microscope (ZEISS, Jena, Germany).
3. Results
The mean grain size and tensile mechanical properties of the materials under study in the CG and UFG states are presented in
Table 4.
The fatigue test results show that the nanostructuring of materials has an ambiguous effect on the total number of loading cycles to failure (life) of the samples.
Table 5 presents selected data on the life of the samples from the CG and UFG materials under study, tested at equal values of the cycle stress range ΔP. It can be seen that the life of the samples from the UFG steels with bcc lattices (9MnSi5, C45 and 12Cr-2W-2Ni-0.5Mo) is higher or marginally lower (at high loading cycles) than the life of the samples from the CG steels. The life of the samples from the CG materials with hcp and fcc lattices (Grade 4 Ti, and the alloys Ti-6Al-4V, Mg6Al and ENAW-2024) is, on the contrary, higher than the life of the samples from the UFG materials (
Table 5).
It is known that the total number of loading cycles to failure (life) of the samples or parts, N, includes the number of cycles to fatigue crack initiation, N
in, and the number of cycles for fatigue crack propagation, N
prop (N = N
in + N
prop). As demonstrated in Refs. [
19,
20], the value of N
in depends on the shape and parameters of the stress raiser. Let us examine the dependence of the N
in values on the life of the samples, N, for the samples of the investigated materials with the same parameters of the stress raisers. It can be seen from
Figure 2a that the number of cycles to fatigue crack initiation (N
in), in both the CG and UFG materials, grows with increasing total life of samples (N). Percentage-wise, the N
in quantity amounts to about 20% of the total life of the samples (
Figure 2b), irrespective of the material state and the crystal lattice type.
The analysis of the straight-line portion of the kinetic diagrams of fatigue fracture for the materials under study (
Figure 3) shows that at the same value of the stress intensity coefficient range (
), for most UFG materials under study, the fatigue crack propagation rate (dl/dN) is close to or even lower than that of the CG materials, irrespective of the crystal lattice type of the materials. The only exceptions are the Ti-6Al-4V titanium alloy (
Figure 3e), where the fatigue crack propagation rate in the UFG alloy is marginally higher than in the CG alloy and the ENAW-2024 aluminum alloy, where the above-mentioned difference in the crack propagation rate is observed only at low values of the stress intensity coefficient (
) (
Figure 3h). This is probably related to the fact that for these CG alloys (alongside Ti), the total number of loading cycles to sample failure is larger than for the samples from the UFG alloys (
Table 5).
The straight-line portion of the kinetic diagrams of fatigue fracture is described by the Paris equation [
16]. Our analysis (
Table 6) shows that for the CG and UFG steels with bcc lattices, the values of the coefficients n in the Paris equation are close (for the C45 and 12Cr-2W-2Ni-0.5Mo steels). For the 9MnSi5 steel, the value of the coefficient n is marginally lower in the UFG state than in the CG state. For all the materials under study with hcp and fcc lattices, the coefficient n is much lower in the UFG state than in the CG state (
Table 6). The low value of the coefficient n in the UFG state, in comparison to the CG state, indicates that the UFG structure is less sensitive to cyclic overload [
2,
20,
21,
22].
Let us analyze these results and establish the fatigue fracture mechanisms of the materials under study in the central part of the fracture surfaces from the perspective of the crystal lattice type.
In the zone of fatigue crack propagation in the 9MnSi5 steel, irrespective of the steel’s state, microrelief consists predominantly of crystallographically oriented fragments, with some places showing fatigue striations and secondary cracks located parallel to them (
Figure 4a,b). On the fracture surface of the steel with a CG structure (
Figure 4a), the striations are clearly visible. In the ECAP-processed steel (
Figure 4b) the striations are not clearly seen, but secondary cracks are clearly visible. The microreliefs of the fatigue fracture surfaces of the C45 steel in the CG and UFG states are similar, and consequently, their fracture mechanisms are similar, as well. In both cases, within the zone of fatigue crack propagation ductile fatigue striations and secondary cracks are visible (
Figure 4c,d). As the fatigue crack grows in the CG 12Cr-2W-2Ni-0.5Mo steel, flat regions where ductile fatigue striations are visible on the surface at a large magnification occupy an ever larger area (
Figure 4e). During the fracture of the samples from the UFG 12Cr-2W-2Ni-0.5Mo steel, in the central part of the fatigue zone there are visible ductile fatigue striations and secondary cracks parallel to the crack propagation front (
Figure 4f).
On the fatigue fracture surfaces of CG Ti, one can observe flat fragments consisting of transcrystalline cleavage-like facets with sizes approximately coinciding with the grain size of CG Ti [
23]. On the surface of the facets there are clearly visible fatigue striations and secondary cracks (
Figure 4g). The fatigue fracture of UFG Ti, at all the stages of crack propagation, is characterized by the formation of a fine microrelief; ductile fatigue striations and secondary cracks are visible (
Figure 4h). The fatigue fracture surface microrelief of the CG Ti-6Al-4V titanium alloy can be characterized as “scaly” or banded [
24]. Such microreliefs are typical for the region of fatigue crack propagation in Ti-6Al-4V. On the surface of large scales, one can observe fatigue striations and secondary cracks (
Figure 4i). The “scaly” microrelief is also preserved on the fracture surface of the UFG Ti-6Al-4V alloy. However, the scales are small; fatigue striations and secondary cracks are visible (
Figure 4j). Within the fatigue zone, the fracture surface of the as-annealed CG Mg6Al alloy consists of transcrystalline cleavage-like facets (
Figure 4k) with sizes approximately coinciding with the alloy’s grain size. The morphology of the facets is represented by equidistant, parallel tubes that have the same orientations within a facet. The formation of such a microrelief is accounted for by the formation of elongated pores in the form of tubes at the intersections of slip bands, and a subsequent rupture of the bridges between them [
25]. The microrelief of the fatigue fracture surfaces of the Mg6Al alloy after ECAP processing is also characterized by a tubular morphology, but the tubes are fragmented (
Figure 4l).
On the fatigue fracture surface of the CG Fe-0.02C-18Cr-8Ni austenitic steel, ductile fatigue striations and secondary cracks are visible (
Figure 4m). On the fracture surface of the UFG steel, the quantity of secondary cracks increases manifold (
Figure 4n). In austenitic steels, this may be attributed to the presence of carbide particles that, on the whole, have a negative effect on low-cycle fatigue resistance, leading to the formation of cracks and voids [
26]. The fracture surface microrelief of the CG ENAW-2024 aluminum alloy may be characterized as a cyclic cleavage with tongues and steps clearly oriented along the crystallographic planes. There are areas with a dimple microrelief (
Figure 4o). On the fatigue fracture surface of the UFG alloy, ductile striations alternate with the cyclic cleavage regions and dimple microrelief (
Figure 4p).
Thus, it can be seen that the fatigue fracture surface microreliefs of the CG and UFG steels with bcc lattices are similar. Apparently, the fatigue fracture mechanisms of the CG and UFG steels are identical. In Ti and the Ti-6Al-4V titanium alloy (hcp lattice materials), strong refinement of the fragmental microrelief forming on the fracture surface of the UFG materials is observed, in comparison to CG materials (except the Mg alloy, where the ECAP-processed UFG state was not achieved); however, the fracture mechanisms of the fragments of the UFG and CG materials are similar. The fatigue fracture mechanisms of the CG and UFG materials with fcc lattices are also similar, although it is noteworthy that the fracture of the UFG materials is accompanied by the formation of many secondary cracks. The last remark is characteristic of the materials with bcc and hcp lattices, as well.