2015 Volume 55 Issue 1 Pages 278-284
Oxide scale formed on the strip surface during hot rolling has posed a serious obstacle to ensure a defect-free surface of steel products in an ecologically friendly way. Recently, an influential idea is that the tertiary oxide scale with a tailored texture can be expected to enhance surface quality and tribolgoical properties during particular lubrication. In this study, texture evolutions of magnetite (Fe3O4) and hematite (α-Fe2O3) in deformed oxide layers formed on a hot-rolled microalloyed steel were investigated by electron back-scattering diffraction (EBSD). Fe3O4 develops a strong θ fibre parallel to the oxide growth, and α-Fe2O3 has a dominant {0001}<1010> texture component. This could be explained by surface energy minimisation during oxides growth and transformation between two oxides, which can also be affected by propagation of cracks along high angle grain boundaries in Fe3O4. Our data further demonstrate that a high thickness reduction (>28%) can reduce α-Fe2O3 wedging through Fe3O4 cracks, and tailoring Fe3O4 texture to {111} components can prevent the α-Fe2O3 growth titling 54.76° from the <001> crystal direction of Fe3O4. As such, these means can dramatically alleviate disturbance from ‘red scale’ (α-Fe2O3) during high-temperature steel processing.
Oxidation is one of the most pervasive problems in almost all metals and their alloys processing at elevated temperatures.1,2,3,4) Specifically, oxide scale formed in hot strip rolling has posed a serious obstacle, and thereby causes the surface quality of final products to deteriorate.5,6,7,8) In most cases, many different surface properties of metals and alloys can dominate surface characteristics of two solids in contact and thereby influence tribological performance.9) These surface properties include surface energy, crystallographic orientation, grain boundaries, texturing of surface and crystal structure.10,11) For instance, some results demonstrate that Fe3O4 (100) plane is a very rich and complex system, and its surface structure depends critically on surface preparation in the magnetoelectronics field.12) It is therefore recently proposed that a tight oxide layer (~75% magnetite, Fe3O4)13,14) with controlled texture, such as Fe3O4 (100), can be possible to enhance tribological properties during nanoparticles-additive lubrication, and further improve surface quality of steels.9,10) This is usually associated with the deformation behaviour of oxide scale during hot rolling.15,16) However, characterisation of crystallographic texture evolution in deformed oxide layers has remained elusive.
Microalloyed low carbon steels generally evolve multilayered oxide scale which consists of two (below 570°C) or three (above 570°C) oxide sub-layers having distinct mechanical properties and growing by different mechanisms.2,3) Correspondingly, wustite (Fe1–xO, with 1–x ranging from 0.83 to 0.95) dominates at high temperatures (normally > 650°C for steels), whereas Fe3O4 prevails at low temperatures, and hematite (Fe2O3) remains low in either case.17,18,19) The proportion of the three phases within oxide scale at high-temperature oxidation depends greatly on the equilibrium transformation of Fe1–xO or Fe3O4 precipitation which is related to heat treatment, atmospheric conditions and alloying elements in the steel.20,21) The growth and microstructure of oxide scale during the continuous cooling from different temperatures and at different cooling rates have been investigated,3,22,23) and it was found that the oxide scale increased in thickness at a lower cooling rate. The microstructural development of oxide scale was observed with Fe3O4 precipitates in the Fe1–xO layer at cooling rates of 10 to 60°C/min.3) The Fe3O4 and iron eutectoid layer formed near the Fe3O4 layer, while a Fe1–xO layer containing Fe3O4 precipitates was detected at the inner region when the cooling rate was reduced to 5°C/min.22) Various explanations for this phase evolution have been proposed, based on an investigation on the isothermal decomposition of the thermally grown Fe1–xO.20,21,24)
Substantial studies on the texture in oxide phases and their orientation relationship (OR) is now being investigated via electron backscatter diffraction (EBSD).17,18,19,25,26,27,28,29) Cubic Fe3O4 and Fe1–xO share a strong <100> texture and a cube-cube OR in undeformed oxide scale whatever the steel substrate.17,18,19) In contrast, the deformed oxide scale develops a pronounced <100> fibre component using plane strain compression.25,26) Alternatively, a noticeable microstructure for the precipitation of Fe3O4 at either high or low temperature is another product of the Fe3O4 layer (called the magnetite seam) formed at the Fe1–xO/steel interface.22,30) Several attempts were made to identify the magnetite seam which was associated with the adherence properties of oxide scale.31,32) In a recent study the orientation of the seam was reported {110}Fe//{100}Fe3O4, <110>Fe//<100>Fe3O4 in the case of transforming by continuous cooling from 400°C.31) Also, their lattice strain was calculated as 4%. By contrast, the Fe/FeO OR was {100}Fe//{110}FeO, <110>Fe//<110>FeO with a lattice strain of 25%. It is suggested that Fe3O4 has better coherency with the steel substrate than Fe1–xO in terms of the conformity of lattice sites of iron oxide with that of the substrate, but these relationships are also limited to the oxide scale thicker than 8 μm.33) For the very thin oxide scale, a Fe/FeO OR was {100}Fe//{100}FeO, <100>Fe//<110>FeO with a lattice strain of 6%.32) Nevertheless, only limited studies on trigonal α-Fe2O3 indicate its grains grow at an angle of 60° from the <100> crystal direction of Fe3O4.17,18)
However, ambiguities still remain what kind of texture in deformed Fe3O4 and α-Fe2O3 will be developed during hot rolling, and what its evolution will be. Little is also currently known concerning grain boundary characters of cubic Fe3O4 and trigonal α-Fe2O3. An attempt will be made in this study to address this deficiency by investigating the texture evolution and grain boundary characters of Fe3O4 and α-Fe2O3 in deformed oxide layers after hot rolling-accelerated cooling (HR-AC) tests.
The material used was a Nb-V-Ti microalloyed low carbon steel in the form of a sheet (400 × 100 × 3 mm3). Its chemical compositions are listed in Table 1. These sheets were then ground using SiC papers of 2400 mesh to a surface finish of 0.5 μm, and cleaned in a solution of acetone. HR-AC experiments were immediately performed on a 2-high Hille 100 experimental mill with 225 mm diameter rolls, combined with a newly developed accelerated cooling system. Full details of the experimental instruments can be found elsewhere.15,34) The following procedure was carried out for every HR-AC test. Each sample was reheated to 900°C at a rate of 1.7°C/s under a high purity inert gas atmosphere, and held for 15 min to ensure a uniform temperature and homogenise the austenite grains. The reheated sample was then given a single rolling thickness reduction (TR, 10%–28%) at a rolling speed of 0.3 m/s without any lubrication followed by an accelerated cooling with the cooling rates (CRs) of 10–28°C/s. Finally, the HR-AC samples were air cooling to obtain the tertiary oxide scale at room temperature. This presumes, of course, the grown oxide layers are subject to the same deformation ratios as the corresponding hot-rolled steel substrate.
| Elements | C | Si | Mn | P | Cr | S | Al | N | Nb +V+ Ti | Fe |
|---|---|---|---|---|---|---|---|---|---|---|
| wt.% | 0.1 | 0.15 | 1.61 | 0.014 | 0.21 | 0.002 | 0.034 | 0.003 | 0.016–0.041 | Bal. |
Sample preparation of oxide scale using conventional polishing techniques is more complicated and time-consuming than that of a single-phase matrix because the oxide phases and the steel substrate need to polish at different rates.35,36) In this study, a Leica EM triple ion beam cutter (TIC020) system was used for precise processing of oxidised samples for subsequent examination with electron microscopes. Samples were cut from the centre of the hot-rolled sheet along the plane of the rolling direction (RD) and normal direction (ND). After gold deposition, the edges for cross sectional analysis were ground by SiC papers with 2000 mesh, and then ion-milled at 6 kV for 5 h using a TIC020 system. Microstructure characterisation was studied using a JEOL JSM 7001F Schottky field emission gun (FEG) scanning electron microscope (SEM) with a Nordlys-II (S) EBSD detector, operated at an acceleration voltage of 15 kV, a probe current of around 2–5 nA, a working distance of 15 mm, and a step size of 0.125 μm.
Post-processing of acquired dataset was carried out using Channel 5 software, where both texture and misorientation data were extracted from the EBSD maps. First for orientation noise reduction, an angular resolution for the grain reconstruction was maintained at a constant value of 2°. Consequently, 2° ≤ θ < 15° misorientations are defined as low-angle grain boundaries (LAGBs), whereas, the high-angle grain boundaries (HAGBs) are θ ≥ 15°. Secondly, phase identification is based on the structural data related to the crystal symmetry and lattice parameters of the suspected phases. In case of the oxide scale, each phase has a different space group (α-Fe, Im3m; Fe1–xO, Fm3m; Fe3O4, Fd3m; α-Fe2O3, R3c), and different lattice parameters. α-Fe, Fe1–xO and Fe3O4 in cubic symmetry with lattice parameters (a) of 0.287 nm, 0.431 nm, and 0.840 nm respectively, and α-Fe2O3 in a trigonal structure with a = 0.504 nm and b = 1.377 nm.17,25,29) The EBSD phase maps were finally divided into three subsets comprising α-Fe, Fe3O4 and α-Fe2O3 for texture based analysis. Orientation distributions of the three subsets were calculated from the data on individual grain orientations collected. The grain orientation g = (ϕ1, Φ, ϕ2) is expressed by the three Euler angles in Bunge notation.36,37) α-Fe and Fe3O4 have a cubic structure and their orientation distribution function (ODF) sections are depicted using the ϕ2 = 0° and 45°, for trigonal α-Fe2O3 the ODF sections with ϕ2 = 0° and 30° are used, in contrast.36) ODF sections were calculated using the discrete binning method with a bin size of 5° and Gaussian smoothing.
The microstructure of the as-received steel substrate was analysed prior to HR-AC experiments. Figure 1 shows a typical SEM image of the etched microalloyed steel using 2% nital solution for 20 s. The microstructure of steel substrate consists of 94% fine-grained polygonal α-Fe and 6% pearlite. The grain size of polygonal α-Fe was determined as 3 μm by using circular intercept method according to ASTM: E112-10. As such, the steel substrate in EBSD based analysis is confined to single α-Fe phase due to only a small amount of pearlite in this material.
Surface morphology of the as-received microalloyed steel substrate.
EBSD phase maps in Fig. 2 indicate the distribution of oxide phases formed on the steel substrate subjected to different TRs and CRs. The oxide scale consists of a two-layered microstructure with a thin outer layer of α-Fe2O3 and an inner duplex Fe3O4 layer. The retained Fe1–xO and eutectoid α-Fe disperse over the Fe3O4 matrix. α-Fe2O3 near the surface gradually penetrates into the cracks within the oxide scale shown in Figs. 2(a) and 2(b). Its distribution in the Fe3O4 matrix reduces dramatically with the decrease of pores or cracks when the TR is high (28% in Fig. 2(c)). It is inferred that an increase of TR can suppress the transformation of Fe3O4 to α-Fe2O3 in the oxide layers after HR-AC process. This is due partly to a refining of Fe3O4 grain size (2–3 μm) at a high TR of 28% (Fig. 2(c)). A large amount of fine granular Fe3O4 grains with relatively dense arrangement can alleviate transformation of Fe3O4 to α-Fe2O3 by oxidation of Fe2+ to Fe3+.
EBSD phase maps for α-Fe, Fe1–xO, Fe3O4 and α-Fe2O3 of the samples with the TRs and CRs of (a) 10%, 10°C/s, (b) 13%, 23°C/s and (c) 28%, 28°C/s.
The EBSD boundary maps in Fig. 3 illustrate a gradual and uniform collapse of the columnar boundary structure. The relatively high TR and CR of 28%, 28°C/s produced large amount of granular Fe3O4 grains (2–3 μm in size). Indeed, the grain-refined Fe3O4 could be affected by Fe1–xO decomposition of hot-rolled samples cooling from high temperature.20,21,24,38) When the similar TRs of 10%, there are much more LAGBs occurred at the low CR of 10°C/s in Fig. 3(a). With the low TR of 13%, Fe3O4 contains very few LAGBs at the CR of 23°C/s (Fig. 3(b)), whereas most of LAGBs were found in the sample with the high TR (Fig. 3(c)). This suggests that a low CR and high TR could enhance the presence of LAGBs. These LAGBs are generally believed to suppress the propagation of cracking during hot rolling. It is further inferred that the coupled effect of TR and CR could be imposed on Fe1–xO decomposition into Fe3O4 below the eutectoid point of 570°C.20,21)
EBSD grain boundary maps of the samples after HR-AC tests at the TRs and CRs of (a) 10%, 10°C/s, (b) 13%, 23°C/s and (c) 28%, 28°C/s. Grey = LAGBs, black = HAGBs.
To further quantity the HAGBs between emerging α-Fe2O3 and Fe3O4, Fig. 4 provides the distributions of misorientation axes/angel (≥15°) of the sample after a TR of 10% and CR of 10°C/s. The distribution of misorientations axes for Fe3O4 in angle ranges has been processed to give a high proportion of coincident site lattice (CSL) boundaries. The peaks in the 57–63° section (Fig. 4(a)) correspond to 60°/<111> (Σ3) misorientation marked in Fig. 4(c). Misorientation peaks occur in α-Fe2O3 (Fig. 4(b)) for axes near <0001> in the angle range of 27–63°, and <1102> in the angle range of 63–83°. The distribution of misorientation angles (Fig. 4(c)) in cubic crystals has a maximum at 45° and a cutoff at 62.8°, whereas trigonal α-Fe2O3 has a maximum at 60° and a cutoff at 95°. For α-Fe2O3 the relatively high densities correspond to 57.42°/<1210> (Σ13b) and 84.78°/<0110> (Σ19c).39) In any case, these grain boundary characters in Fe3O4 and α-Fe2O3 can provide deep insight for the penetration of α-Fe2O3 into Fe3O4 scale during air cooling after hot rolling.
Distributions of misorientation axes for (a) Fe3O4, (b) α-Fe2O3, of the sample with a TR of 10% and a CR of 10°C/s, and (c) corresponding distribution of misorientation angels (≥ 15°).
A typical quantitative evaluation was conducted to obtain the texture evolution in Fe3O4/α-Fe2O3 scale at the different TRs and CRs. The representation of pole figures (PF) for Fe3O4/α-Fe2O3 scale at a TR of 13% and a CR of 23°C/s has been given in Fig. 5. Fe3O4 develops the strong {100} and {110} fibres and a relatively weak {111} texture component (Fig. 5(a)). The {0001} basal planes of α-Fe2O3 are approximately parallel to the interface with Fe3O4, thus building the fibre axis along the surface normal direction (Fig. 5(b)). The {1120} and {1010} prismatic planes lie nearly perpendicular to the magnetite surface. The results are slightly different from the previous investigation where the {1120} and {1010} fibres texture develop instead.39) This could be due to the oxide scale without deformation in previous studies.39)
(100) and (111) pole figures (PF) for (a) Fe3O4 and (b) (0001), (1010) and (1120) PF for α-Fe2O3 at a TR of 13% and a CR of 23°C/s.
Figure 6 shows texture development of Fe3O4 and α-Fe2O3 in deformed oxide layers and their intensity distributions along associated fibres or texture components. Fe3O4 has a cubic structure and its ODF sections are depicted using the ϕ2 = 0° and 45° (Fig. 6(b)) in terms of the Bunge system.37) In contrast, for trigonal α-Fe2O3 the ODF sections with ϕ2 = 0° and 30° (Fig. 6(d)) are used.36) θ fibre develops in Fe3O4 superimpose on ϕ2 = 0° section at ϕ = 0° with the rotations of <100>//ND.17,19,25,29) Figure 6(a) shows the intensity distribution of the θ fibre in Fe3O4 on the samples subjected to various deformation conditions. <1010> fibre component in α-Fe2O3 (Fig. 6(c)) lies on ϕ2 = 30° section and corresponds to orientations along Φ = 90°.
Development of texture intensity f(g) along the (a) θ fibre of Fe3O4, (c) <1010> fibre of α-Fe2O3, of the samples with different TRs and CRs, and ODF sections for (b) Fe3O4 and (d) α-Fe2O3 at a TR of 28% and a CR of 28°C/s.
As seen in Fig. 6(b), Fe3O4 in deformed oxide layers holds a strong θ fibre parallel to the oxide growth, where pronounced orientations develop between {001}<100> cube and {001}<110> rotated cube components. With the increase of TRs and CRs (Fig. 6(a)), the cube orientations slightly shift toward the <210> direction, and finally lead to a relatively strong {100}<210> component with the intensity of f(g) = 6.7. θ fibre is also found in the undeformed Fe3O4 left at 650 and 950°C after reheating at 1050°C.25,26) Thus, the presentence of {100}<001> cube and {001}<110> rotated cube components can be attributed to grain growth during high temperature oxidation. Nevertheless, the {100}<210> texture can dominate during subsequent deformation by external force. Moreover, it cannot be seen any texture components along {111} slip planes occurred in the most face-centred cubic (fcc) materials.
Compared to the cubic Fe3O4, the trigonal α-Fe2O3 in deformed oxide layers has similar texture components with the nearly same strength under various deformation conditions. The pronounced orientations of α-Fe2O3 (Fig. 6(d)) comprise a fibre running along {0001} basal plane which includes a dominant {0001}<1010> and relatively weak {0001}<2110> texture components. Particularly, the maximum texture intensity lies on the {0001}<1010> orientation (f(g) = 14.4 in Fig. 6(c)) of the sample at a higher TR of 28% and CR of 28°C/s. The {0001} basal plane of α-Fe2O3 is crystallographic plane having close-packed oxygen anions, thus oxide growth and deformation tend to occur across close-packed oxygen planes.17,40) Nonetheless, α-Fe2O3 in a trigonal structure generally has independent slip systems, namely {0001}<2110>, {1210}<1010>, {1210}<1011>, {1102}<0111>, and {1010}<0111>.36,39,41) In this case, the rolling texture development of α-Fe2O3 may attribute to the effect oxygen content through changes of the critical resolved shear stress of the various slip systems.36,41)
The results presents here demonstrate that Fe3O4 develops a strong θ fibre parallel to the oxide growth, and α-Fe2O3 carries a fibre running along {0001} basal plane. This finding that texture evolution of Fe3O4 and α-Fe2O3 positively related to TRs is similar to those two previous studies.25,26) Here it is possible to find out the cause of what is happening. The driving energy for the sharp texture formation observed in Fe3O4 and α-Fe2O3 can attribute to surface energy minimisation when the steels suffering high temperature oxidation during hot rolling. For cubic Fe3O4 with a spinel structure, the surface energy increases in order of {100} (1.5 J m–2), {110} (1.8 J m–2) and {111} (2.2 J m–2).17,18) Hence, this results in the main θ fibre running through {001}<100> cube and {001}<110> rotated cube components (Fig. 6). Indeed, previous results also reveal that the (001) planes of Fe3O4 depends critically on sample preparation and heat treatment.10,42) In case of α-Fe2O3 with a strong preferential crystallographic orientation in the normal direction of the (0001) basal plane, the texture strength is thought to be driven by the low surface energy (1.52 J m–2) of the (0001) plane of α-Fe2O3.41) The rolling texture of α-Fe2O3 has a strong {0001}<1010> component (Fig. 6) which is a direct consequence of the favoured basal plane slip.
Transformation of Fe3O4 to α-Fe2O3 by oxidation of Fe2+ to Fe3+ can be another closely related process responsible for the observed texture evolution in deformed oxide layers. Fe2+ ions oxidised to Fe3+ during high-temperature short-time processing, such as the tertiary oxide scale,3,11,13) are dependent primarily on crystallographic similarities of Fe3O4 and α-Fe2O3. (111) planes of Fe3O4 and (0001) planes of α-Fe2O3 are both parallel and perpendicular to the three-fold axes. For a fcc crystal interfaced to a trigonal close packed (hcp) crystal, the preferred orientation relationships are given by (111)fcc//(0001)hcp and (110)fcc,//(1010)hcp, where the close-packed planes and directions are matched across the interface.40,41) Given that the formed Fe3O4 after thermodynamically unstable Fe1–xO has <001> fibre texture (Fig. 6), and also the interplanar angle between (111) and (001) planes is 54.76° in cubic structure, α-Fe2O3 grains must grow at an angle of 54.76° from the <001> crystal direction of Fe3O4. Thus, Fe2+ ions resulting from diffusion at Fe3O4 boundaries react with oxygen and oxidise as thin platelets of α-Fe2O3 with c-axes perpendicular to the oxide growth and the a-axes in the plane of oxide growth ({0001}<1010> texture component). These results confirm that if α-Fe2O3 grows on a (100) Fe3O4, its grains will be titled at an angle of 54.76° to the Fe3O4 matrix, whereas for a (111) Fe3O4, the epitaxial growth of α-Fe2O3 (0001) will occur. Similar result has been recently reported for other single crystal oxide scales.41)
In addition to the texture developed in undeformed oxide scale, the variation of crystallographic planes is dominant in Fe3O4 and α-Fe2O3 as the different ORs of steel substrate.17,19,29) For most polycrystalline substrates, a noticeable feature in Fe3O4 is the evolution of crystallographic direction, whereas the texture evolution of α-Fe2O3 varies in the crystallographic planes.17,29) This is a direct consequence of the minimisation of surface energy and lattice strains from their mismatch. In case of Fe3O4, the strong textures generally develop on the planes of {100} whatever in undeformed or deformed oxide scale.25,29) The difference is that the deformed Fe3O4 holds the texture development shifting from the <100>, <110> to <210> directions in Figs. 5(a) and 6(b). Nevertheless, α-Fe2O3 is totally another story before and after deformation of oxide scale. The undeformed α-Fe2O3 formed at low-temperature oxidation has the {1120} and {1010} fibre textures,29,41) whereas the rolling texture of α-Fe2O3 has a strong {0001}<1010> component (Fig. 5). This infers that the favoured basal plane slip dominates the deformed α-Fe2O3 compared to the undeformed oxide scale.
Lastly, the formation and propagation of cracks along the Fe3O4 grain boundaries create paths for easy air penetration, and therewith allow further oxidation (Fig. 2(a)). It is believed that low-angle and low-ΣCSL boundaries in microstructure offer obstacles to the propagation of cracks, as they minimise the solute effects and reduce the interaction between the interfaces and glissile dislocation.43) In this case, the oxide scale is easy to crack in presence of Σ13b and Σ19c in α-Fe2O3 compared to Fe3O4 with Σ3 (Fig. 4(c)). Therefore, two means can be considered for the decrease and further easy removal of α-Fe2O3. One is that a desired TR (>28%) can reduce cracks along the Fe3O4 grain boundaries and thereby prevent further oxidation to α-Fe2O3. Alternatively, tailoring the texture of initial Fe3O4 to {111} components using particular lubrication can avoid α-Fe2O3 growth titling 54.76° to the (100) Fe3O4 matrix. This is practically important to dramatically alleviate disturbance from ‘red scale’ (α-Fe2O3) during hot rolling, when Fe3O4 is oxidised to red powdery α-Fe2O3 during air cooling.6,8)
In this study microstructure and texture evolution of Fe3O4/α-Fe2O3 scale form on a Nb-V-Ti steel subjected to the TRs of 10%–28% and CRs of 10–28°C/s were systematically analysed. The following conclusions can be drawn.
(1) The oxide scale consists of a thin α-Fe2O3 outer layer and an inner duplex Fe3O4 layer. Grain refinement of Fe3O4 (2–3 μm) at a TR of 28% can suppress the penetration of α-Fe2O3 into the oxide scale.
(2) The oxide scale comprises a high proportion of LAGBs at low CR and high TR. These LAGBs are generally believed to suppress the propagation of cracking during hot rolling. For CSL boundaries, Fe3O4 on the samples at a TR of 10% and CR of 10°C/s develops 60°/<111> (Σ3), and α-Fe2O3 for 57.42°/<1210> (Σ13b) and 84.78°/<0110> (Σ19c).
(3) Fe3O4 carries strong θ fibre parallel to the oxide growth, which includes {100}<001> and {001}<110> texture components, because of surface energy minimisation on {100} planes of Fe3O4. α-Fe2O3 develops {0001}<1010> component as the favoured basal plane slip. The α-Fe2O3 growth on θ fibre Fe3O4 is 54.76° titled from the <001> crystal direction of Fe3O4.
(4) Two means are proposed for easy removal of α-Fe2O3. One is that LAGBs can act as obstacles to cracks propagation further to prevent α-Fe2O3 penetration. Another, Fe3O4 with {111} texture components can used to form uniform α-Fe2O3. The understanding we have gained as to how texture and grain boundary take place could aid in the choice and design of desired oxide scale available through high-temperature steel processing.
We are grateful to J. Abbott, L. Hao and S. Rodd for their assistance in the HR-AC experiments. The authors also acknowledge use of facilities within the UOW Electron Microscopy Centre.
