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DOI: 10.1002/adem.201000312

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Structure and Proper ties of Nanograined Fe­C Alloys after Severe Plastic Deformation**
By Boris B. Straumal,* Sergei V. Dobatkin, Alexei O. Rodin, Svetlana G. Protasova, Andrei A. Mazilkin, Dagmar Goll and Brigitte Baretzky

The microstructure and properties of several Fe­C alloys are studied in a) the as-cast state, b) after a long annealing time at 725 8C, and, c) after high-pressure torsion (HPT) The grain size after HPT is in ¨ the nanometer range. Only Fe3C (cementite) and a-Fe remain in the alloys after HPT. Less stable Hagg carbide and retained austenite disappear after HPT, and phase composition closely approaches the equilibrium corresponding to the HPT temperature and pressure. This HPT behavior differs from that associated with ball-milling, which can lead to the formation of metastable or amorphous phases. Therefore, severe plastic deformation opens the way to produce materials with very stable phase structure and thus ensures the stable properties of nanograined steel during its life-time.
Some years ago, we observed that severe plastic deformation (SPD) leads to the extremely quick decomposition of the supersaturated solid solution in Al-based alloys.[1,2] During SPD, the phase composition closely approached the equilibrium state at SPD pressure and temperature. We concluded that SPD can be considered as ``hot deformation at room temperature'', i.e., a balance between deformation-induced grain refinement and deformation-accelerated formation of equilibrium phases. Quite recently it has been discovered in in situ synchrotron experiments that enormous amount of crystal defects (for all, vacancies) forms during SPD.[3] They can be responsible for the accelerated diffusion needed for the transformation of metastable into stable phases during SPD. Therefore, we decided to investigate in this work the structural changes in Fe­C alloys during high pressure torsion (HPT) for two reasons: first, due to the enormous importance of Fe­C system, and, second, due to the miraculous phenomenon of ``cementite disappearance'' in steels during SPD, which has to be clarified.[4­12]

[*] Prof. B. B. Straumal, Dr. B. Baretzky Karlsruhe Institute for Technology (KIT), Institute for Nanotechnology, Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, (Germany) E-mail: straumal@mf.mpg.de Prof. B. B. Straumal, Dr. S. G. Protasova, Dr. A. A. Mazilkin Institute of Solid State Physics, Russian Academy of Sciences, Chernogolovka, 142432 (Russia) Prof. B. B. Straumal ¨ Max-Planck-Institut fur Metallforschung Heisenbergstraúe 3, 70569 Stuttgart, (Germany) Prof. S. V. Dobatkin A.A. Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Leninsky prosp. 49, 119991 Moscow, (Russia) Dr. A. O. Rodin National Research University MISIS, Leninsky prosp. 4, 119991 Moscow, (Russia) Dr. D. Goll Hochschule Aalen, Beethovenstraúe 1, D-73430 Aalen, (Germany) [**] The authors thank the Russian Foundation for Basic Research (contracts 09-03-00784, 09-03-92481 and 09-08-90406), the Israel Ministry of Science (contract 3-5790), and Ukrainian Fundamental Research State Fund (contract X28.7049). They also greatly appreciate Dr. A. Nekrasov for stimulating discussions of magnetic properties. Prof. R. Valiev and Dr. Yu. Ivanisenko are heartily acknowledged for the fruitful discussions on severe plastic deformation.

Results In Figure 1, the concentration dependence of the lattice parameter in a-Fe (ferrite) after HPT is shown compared with lattice parameter of non-deformed pure a-Fe.[13] After HPT the lattice parameter of the a-Fe solid solution does not depend on the total carbon concentration in the Fe­C alloys. The lattice parameter is about 0.28667 nm. The solubility of carbon in a-Fe at 680 8C is 0.018 wt.% C.[14] The addition of 0.018 wt.% C to pure a-Fe increases the lattice parameter of a-Fe by 7 á 10þ5 nm.[15] In our case, the lattice parameter after HPT is about 4 á 10þ5 nm higher than that of pure a-Fe. This

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Fig. 1. Dependence of lattice parameter a in a-Fe (ferrite) after HPT on carbon concentration. The lattice parameter of non-deformed pure Fe is taken from ref.[13].

means that the carbon concentration in a-Fe after HPT is about 0.01 wt.% C and does not exceed the carbon maximum solubility in the ferrite (about 0.02 wt.% at 740 8C). Within the framework of experimental error, this is more or less equal to the solubility at room temperature.[14] In other words, HPT does not lead to the formation of supersaturated carbon solid solution in a-Fe. The splitting of the (200) x-rays diffraction line typical for tetragonal martensite lattice was not observed either before or after HPT. This means that both hypoetectoid and hypereutectoid FeþC alloys do not contain martensite. In Figure 2 the bright field TEM image of the as-cast Fe­1.7% C alloy and respective selected area electron diffraction (SAED) pattern are shown. Transformed austenite grains contained the troostite lamellar colonies with interlamellar spacing of 100­400 nm. The boundary between two troostite colonies is visible in Figure 2a. This boundary contains the cementite layer. A white iron-enriched zone is visible along this GB cementite layer. Inside the bright ferrite grains, the fine cementite lamellae appear dark. The presence of ferrite and cementite in the structure is supported by the SAED pattern (Fig. 2b). In Figure 3 the bright-field TEM image of another location in the coarse grained as-cast Fe­1.7 wt.%C alloy is shown together with the respective SAED pattern. This illustrates the presence of retained austenite and a small amount of Hagg carbide in the as-cast alloys. Reflections of ¨ cementite were also detected in this SAED pattern but are not indicated for the reasons of clarity. TEM of the as-cast alloys with 0.25, 0.45, 0.60, 1.3, and 1.5 wt.% C revealed the presence of ferrite, cementite and retained austenite in all studied alloys. Only the amount of these phases is different in various alloys, in conformity with Fe­C phase diagram.[16] The amount of Hagg carbide in the as-cast alloys is very small; ¨ the respective diffraction spots are not present in all studied locations. According to the LM and TEM, the grain size of ferrite (in hypoeutectoid alloys), retained and transformed austenite (in all alloys) scattered from 200 to 700 mm. TEM and LM also do not show any presence of martensite in the as-cast state of material. Neither split (200) spots (typical for tetragonal martensite lattice) in the SAED patterns nor typical martensite grain morphology (needles, plates) were discov-

Fig. 2. a) Bright-field TEM image of the coarse grained as-cast Fe­1.7%C alloy, and, b) corresponding electron diffraction pattern. Fine cementite lamellae appear dark inside of bright ferrite grains.

ered in LM. Electron diffraction, similar to XRD did not show any presence of graphite in all samples studied. Figure 4 shows the LM microstructures of the FeþC alloys with 0.3, 0.6, 1.3 and 1.7 wt.% C after annealing for 950 h at 725 8C (i.e., below the eutectoid temperature). The long annealing time was used in order to produce the equilibrium a-Fe × Fe3C structure. The structures of hypoeutectoid and hypereutectoid alloys became very similar after long annealing. All samples contain very coarse ferrite and cementite grains. Only the amount of cementite increases with increasing carbon content, according to the FeþC phase diagram.[16] Severe plastic deformation by HPT produces the nanometer range grain structure in the FeþC alloys. Figure 5 shows bright- (Fig. 5a) and dark-field (Fig. 5b) TEM micrographs of the Fe­0.3 wt.% C alloy. SAED pattern (Fig. 5c) contains only a-Fe and Fe3C spots. The TEM micrographs and SAED patterns of Fe alloys with 0.45, 0.60,

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Fig. 3. a) Bright-field TEM image of another location in the coarse grained as-cast Fe­1.7% C alloy, and, b) corresponding electron diffraction pattern. Retained austenite ¨ and small amount of Hagg carbide are present.

1.3, 1.5 and 1.7 wt.% C after HPT are very similar to those shown in Figure 5. Both the ferrite grains and cementite particles are visible in the DF image (Fig. 5b) as the reflections of these phases lie closely to each other. Dislocation density is > 1014 mþ2 in all samples. Ferrite grain size after HPT is about 100 nm, which increases slightly with increasing carbon content. Ferrite grains are not equiaxial; they are slightly elongated parallel to the deformation direction. Cementite grain size in all studied alloys is about 30 nm. Their shape is more equiaxial than that of ferrite grains. Fe3C grains are more or less uniformly distributed over the specimen. The spacing between cementite particles decreases with increasing carbon content. It is about 50­100 nm in Fe­0.3 wt.% C alloy and about 10­50 nm in the Fe­1.7 wt.% C alloy. Only two phases, namely a-Fe and Fe3C are present after HPT in all studied alloys. No signs of retained cementite, graphite or other iron carbides are present in the SAED patterns.

In Figure 6a to c the XRD spectra for the Fe­1.7 wt.% C alloy are shown. The initial as-cast alloy (Fig. 6a) contains three phases. First, a strong peak of a-Fe with body-centered cubic (bcc) lattice is present (marked by a filled square). Second, peaks of retained austenite (face-centered cubic, g-Fe) are detected (marked by triangles). Third, the weak and broadened peaks of cementite are also present (marked by arrows). Some very weak peaks (like those at 2u ¼ 42.23, 55.3 and 66.488) can be attributed as diffraction from the Hagg ¨ carbide Fe5C2. Figure 6b contains the data for the Fe­1.7 wt.% C alloy annealed for 950 h at 725 8Cinthe a × Fe3C two-phase region of the FeþC phase diagram. The spectrum contains only two phases, namely ferrite and cementite. No peaks of retained austenite and Hagg carbide are present. Figure 6c ¨ contains the diffraction spectrum for the Fe­1.7 wt.% C alloy after HPT. Similar to the results in Figure 6b, only two phases (ferrite and cementite) remained after HPT in the sample and retained austenite and Hagg carbide disappeared. However, ¨ all peaks are broadened even in comparison with the as-cast state due to the fine grain size (and/or residual stresses) in samples after HPT. XRD did not show any presence of graphite in all studied samples. The lines of ferrite (Heff ¼ 330 kOe[17]), Hagg carbide ¨ x-Fe5C2 with hyperfine fields of 185 ô 3 kOe,[18] cementite Fe3C (Heff ¼ 210 ô 5 kOe[17]) and retained austenite (no hyperfine field)[17] were found in the Mossbauer spectra of ¨ the as-cast alloys (Fig. 7a). The Mossbauer spectra after HPT ¨ contain only the lines of ferrite and cementite Fe3C (Fig. 7c and 7d). Retained austenite and Hagg carbide disappeared ¨ both after long annealing (Fig. 7b) and after HPT (Fig. 7c and d). Other iron carbides like e-Fe2C with hyperfine fields of 170 ô 3, 237 ô 3 and 130 ô 6 kOe,[19,20] Fe4C or Fe6C were not observed either before or after HPT. This fact supports the data from XRD and TEM. The amount of carbides before and after HPT calculated based on Mossbauer measurements ¨ increases linearly with increasing carbon content. However, the amount of carbides in the as-cast Fe­C alloys is much higher than that in the HPT alloys. This indicates decrease of the cementite amount after HPT. The samples after HPT became rather brittle, which made it possible to break them in situ in the Auger spectrometer chamber. The carbon concentration along GB fracture surfaces fluctuates between 10 to 70 wt.% in the layer of about 1 to 3 nm thickness. This reveals the strong grain boundary carbon adsorption, but not the continuous cementite GB layers. Discussion It has been observed recently that after HPT of AlþZn alloys with 10, 20 and 30 wt.% Zn, the Zn- supersaturated (Al) solid solution decomposes and closely approaches the equilibrium state corresponding to the room temperature.[1,2] The decomposition of the supersaturated (Al) solid solution proceeds parallel to the drastic grain refinement. No metastable phases conventionally produced at the AlþZn alloys decomposition (Guinier-Preston I and II zones,

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Fig. 4. Light micrographs of the FeþC alloys with a) 0.25, b) 0.6, c) 1.3, and, d) 1.7 wt.% C annealed at 725 8C for 950 h.

Fig. 5. a) Bright-, and, b) dark-field TEM micrographs of the Fe­0.3 wt.% C alloy after HTP. c) Electron diffraction pattern revealing a-Fe and Fe3C spots.

rhombohedral distorted fcc aR -phase or dis0 torted fcc am -phase) were observed after SPD. We concluded that HPT is nothing else than ``hot'' deformation at room temperature. In other words, the vacancy production during HPT enhances the diffusion to such a high extent, that the equilibrium at room temperature phases appears parallel to the drastic grain refinement. The high vacancy concentrtation during HPT has been revealed recently in situ using the synchrotron radiation.[3] In the case of FeþC alloys, ferrite and graphite are equilibrium phases at room temperature. However, cementite becomes equilibrium at rather low pressures (0.1­0.5 GPa, depending on the temperature).[21­23] In our case the pressure during HPT was much higher, 5 GPa, ensuring the thermodynamic stability of cementite. Other iron carbides become stable above 5 GPa. For example, e-carbide Fe7C3 becomes stable above 5.9 GPa.[23] Therefore, such carbides do not appear in our alloys after HPT. However, it may be the reason why the HPT of the U13 steel containing 1.37 wt.% C at 12 GPa resulted in appearance of e-carbide Fe7C3 and Hagg carbide Fe5C2 at the cost of cemen¨ tite.[9] In early experiments on the severe plastic deformation of steels, Korznikov et al. [24] observed the disappearance of the cementite peaks from XRD spectra and assumed that, similar to ball-milling, SPD leads to the formation of the supersaturated carbon solid solution in a-Fe. Later is was observed that cementite becomes fine but does not completely disappear from the SPD-treated steel.[8,1,13,25] Nevertheless, we ask why the amount of cementite decreases after SPD? Nanograined materials contain a large number of interfaces (grain and interphase boundaries). Up to 10% of atoms can be positioned in these interfaces. In all two- or multicomponent systems the interface segregation takes place. This means that the composition of interfacial layers is generally not equal to the overall composition of the bulk material. Normally, the interfaces are enriched by one of the components. In certain cases thin thermodynamically stable layers of an intergranular phase may form.[26] Conventional experimental methods usually applied for the investigation of bulk phases are generally not able to detect the input of GB segregation or thin GB phases. For example the XRD peaks are caused by rather large areas of coherent scattering. If the particles of a phase become too small, the X-ray peaks

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there is a big difference between interface composition and overall bulk composition. FeþC alloys are a good example of a system with strong GB segregation. The carbon atoms strongly segregate in the ferrite GBs. The concentration of carbon atoms in the ferrite GBs measured in the hypoeutectoid steels lies between 0.2 and 0.5 monolayer.[27] Our AES measurements for the hypereutectoid alloy with 1.7 at.% C also reveal several tens of percent of carbon in the ferrite GBs. It can be seen that cementite particles become very fine, but do not completely disappear from the steel (Fig. 5). On the other hand, carbon also does not form the supersaturated bulk solid solution in ferrite. Its concentration does not exceed the equilibrium bulk solubility limit in the a-Fe (Fig. 1). However, the magnetic data reveal that the overall amount of the bulk cementite phase after HPT becomes lower than the equilibrium phase diagram predicts.[28] The magnetization of the samples after HPT is much higher than the Js values of the as-cast coarse-grained FeþC alloys. This means that the large amount of carbon has ``disappeared'' in the ferrite GBs and cannot form bulk cementite. As a result, the amount of cementite in the nanograined HPT alloys is much lower than in the as-cast or coarse-grained FeþC alloys. This is the reason of the apparent disappearance of bulk cementite after HPT. About one third of bulk cementite disappeared after HPT.[28] The respective carbon amount formed GB segregation layer and does not input into the overall magnetization value. After HPT the grain size in the studied FeþC alloys is about 100 nm. If we suppose that the GB thickness is about 0.5 nm, the total volume of GB layers would be about 1.5 at.%. The magnetization of Fe­1.7 wt.% C alloy after HPT is nearly equal to the magnetization of coarse-grained alloy with 1.2 wt.% C.[28] This means that ferrite/ferrite GBs contain about 0.5 wt.% C. The estimation of the resulting carbon concentration would be about 0.3 carbon monolayers. This value obtained from the magnetic measurements corresponds well both with our Auger data and measurements.[27] Conclusions The following conclusions can be drawn. i) After high pressure torsion (HPT) of FeþC alloys, the non-equilibrium (metastable) phases disappear from the alloys. The phases which are equilibrium at temperature and pressure of HPT appear instead. It is an important difference from ball-milling which can lead, similar to the ion implantation, to the formation of metastable or amorphous phases. Therefore, severe plastic deformation opens the way to produce materials with very stable phase structure and thus ensures the stable properties of nanograined steel during its life-time. ii) HPT leads to the strong grain refinement of FeþC alloys. Carbon actively segregates in the numerous ferrite/ferrite grain boundaries (GB) and triple junctions. The formation of new GB segregation layers during SPD consumes carbon and leads to the apparent disappearance of bulk cementite. Bulk cementite, indeed dissolves, however, not in the bulk ferrite but in ferrite/ferrite GBs.

Fig. 6. X-ray diffraction spectra for b) after annealing at 725 8C for 950 ferrite peak is marked by the square. mark the retained austenite. Arrows

the Fe­1.7 wt.% C alloy: a) initial as-cast alloy. h, and, c) after HPT at 5 GPa, 5 torsions. (110) ¨ Crosses mark the Hagg carbide Fe5C2. Triangles mark the peaks of cementite Fe3C.

broaden and finally ``sink'' in the background. In other words, the amount of bulk phases in the polycrystalline alloy can strongly differ from that predicted by the equilibrium phase diagram if: i) the grain size is in the nanometer range and the portion of atoms, which are positioned not in the bulk but at the interfaces, is high; or, ii) strong interface segregation, i.e.,

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¨ ¨ Fig. 7. Mossbauer spectra of the FeþC alloys. Squares mark the ferrite peaks. Arrows mark the peaks of cementite Fe3C. Crosses mark the Hagg carbide Fe5C2. Triangle mark the retained austenite. a) As-cast Fe­0.6 wt.% C alloy. The area under the spectrum components corresponds to the 10% of Fe3C and Fe5C2 and 5% of retained austenite. b) As-cast Fe­0.6 wt.% C alloy after long anneal. The area under the spectrum components corresponds to the 8% of Fe3C and Fe5C2.Fe­0.6 wt.% C alloy after long annealing. The area under the spectrum components corresponds to the 8% of Fe3C and Fe5C2. c) The Fe­0.45 wt.% C alloy after HPT. The area under the spectrum components corresponds to the 4% of Fe3C. d) The Fe­1.5 wt.% C alloy after HPT. The area under the spectrum components corresponds to the 12% of Fe3C.

Experimental
Both hypo- and hypereutectoid FeþC alloys with carbon concentration of 0.05, 0.10, 0.20, 0.25, 0.45, 0.60, 1.3, 1.5 and 1.7 wt.% were prepared from high purity 5N Fe and C by vacuum induction melting in the form of cylindrical 12 mm diameter ingots. The carbon content was measured by atomic absorption spectroscopy in a Perkin-Elmer spectrometer. 2 mm thick discs were subjected to HPT at room temperature in a Bridgman anvil-type unit under a pressure of 5 GPa and for 5 torsions. Samples for structural and magnetic investigations were cut from the HPT-deformed discs at a distance of 3 mm from the sample center. One set of as-cast samples with 0.25, 0.60, 1.3 and 1.7 wt.% C was additionally annealed for 950 h at 725 8C in order to achieve the equilibrium a × Fe3C structure. Light microscopy (LM) was performed with a Zeiss Axiophot microscope. Transmission electron microscopy (TEM) investigations were carried out on a JEM­4000FX microscope at accelerating voltage of 400 kV. X-ray diffraction (XRD) data were obtained on a Siemens diffractometer (Co Ka radiation). Mossbauer experiments were performed ¨ using a Perseus spectrometer with Co57 source in the rhodium matrix at room temperature in the velocity interval from þ10 to 10 mm à sþ1 with 256 registration channels. This allowed the effective magnetic field to be defined with an error of less than 5 kOe. Spectra were analyzed using the least-square method for the superposition of Lorentz lines corresponding to the absorption of various phases.

Auger-electron spectroscopy (AES) of HPT Fe­1.7 wt.% C alloys was performed using a PHI 680 Auger spectrometer.

Received: October 11, 2010 Final Version: February 10, 2011 Published online: March 10, 2011

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