Columnar grain growth during annealing prior to cold rolling of non-oriented electrical steels

Materials Science & Engineering B 243 (2019) 8–18

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Columnar grain growth during annealing prior to cold rolling of nonoriented electrical steels

T

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E.J. Gutiérrez Castañedaa, , C.N. Palafox Cantúb, A.A. Torres Castillob, A. Salinas Rodríguezc, R. Deaquino Larac, F. Botello Riondac, F. Márquez Torresc, S. García Guillermoc a Catedrático CONACYT – Instituto de Metalurgia de la Universidad Autónoma de San Luis Potosí, Av. Sierra Leona, 550, Lomas 2da Secc., San Luis Potosí, S. L. P. 78210, Mexico b Instituto de Metalurgia de la Universidad Autónoma de San Luis Potosí, Av. Sierra Leona, 550, Lomas 2da Secc., San Luis Potosí, S. L. P. 78210, Mexico c Centro de Investigación y de Estudios Avanzados del Instituto Politécnico Nacional (CINVESTAV-Saltillo), Av. Industria Metalúrgica, 1062, Ramos Arizpe, Coahuila 25000, Mexico

A R T I C LE I N FO

A B S T R A C T

Keywords: Electrical steels Columnar growth JMatPro simulation Phase transformation Crystallographic texture Microstructure

The present research reports the microstructural factors that favor the development of columnar grains during annealing prior to cold rolling (APCR). Results show that decarburization of steel has a strong effect on phase transformations, grain structure and crystallographic texture. Annealing at 850 °C (two-phase field) promotes the development of columnar-grained microstructures with orientations close to the ideal cube texture. Higher annealing temperatures within the intercritical region, results in a mixture of columnar + equiaxial grains (875 °C) or completely equiaxial-grained microstructures (900 °C). Annealing at temperatures within the single ferrite (700 °C, 750 °C) or austenite (950 °C, 1050 °C) phase fields does not promote the development of columnar grains. In these cases, annealing promotes the development of equiaxial grains with rotated cube texture and orientations along the γ-fibre. Changes in texture produced by APCR are consistent with the Bain and Kurdjumov-Sachs correspondence relationships and occur according to oriented nucleation and selective growth mechanisms depending on the characteristics of decarburization.

1. Introduction Grain non-oriented (GNO) electrical steels play an important role in the generation, distribution and consumption of electrical energy, which represents a great concern worldwide in terms of available resources, environmental and economic reasons [1]. Their applications include generators, motors, smaller-sized units within the automotive sector, industrial machines and household appliances [2]. For this reason, these materials are considered as one of the most important among the magnetic materials produced today [1,3]. For these applications, high permeability and low core loss are required; these magnetic properties depend strongly on chemical composition, grain size, residual stresses, secondary particles and crystallographic texture [2,4]. Conventional manufacture of GNO electrical steels from continuous casting involves hot-rolling, cold rolling, continuous or batch annealing and temper rolling [5]. Finally, after stamping from strip blanks, laminations are subjected to a long-term decarburization annealing [6]. Unfortunately, this process cannot obtain the appropriate

microstructure that maximizes the energy efficiency [7]. Much work has been carried out to tailor the magnetic properties of GNO electrical steels. The development of columnar grains is as an efficient method to improve the magnetic quality (less core losses and higher permeability) of these steels, which is related to certain microstructural changes that take place during this growth such as reduction in amount of carbides [8], increase of grain size [9] and development of cube or rotated cube textures [10]. Most of information regarding the columnar grain growth in non-oriented electrical steels involves decarburization annealing after cold rolling [8–10]. In the present work, decarburization of steel and development of columnar grains take place prior to cold rolling. Instead of using a long-term vacuum annealing [11] or a stepped-temperature annealing as recommended by other authors for temper-rolled or cold-rolled steels (to promote recrystallization and phase transformation) [12], a simpler isothermal short-term annealing treatment is suggested to develop columnargrained microstructures. Results of previous works related to the effects of annealing prior to

⁎ Corresponding author at: Institute of Metallurgy of The Autonomous University of San Luis Potosi (UASLP), Department of Metallurgy and Materials Engineering, 550 Sierra Leona Avenue, Lomas 2nd Section, San Luis Potosi, San Luis Potosi 78210, Mexico. E-mail address: [email protected] (E.J. Gutiérrez Castañeda).

https://doi.org/10.1016/j.mseb.2019.03.016 Received 6 February 2018; Received in revised form 13 January 2019; Accepted 21 March 2019 0921-5107/ © 2019 Published by Elsevier B.V.

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finish with colloidal silica (0.05 μm). Texture analysis were carried out from color-coded maps of the inverse pole figure (IPF [0 0 1]), and determination of the orientation distribution function (ODF). The analyzed area for each orientation map was 800 μm × 1250 μm, and the ODF was calculated by the harmonic series expansion method using the TSL-OIM 3.5 commercial software.

cold rolling (APCR) on the microstructure and magnetic properties of non-oriented electrical steels show that columnar-grained microstructures lead to an improvement of the magnetic behavior of the final product (obtained after cold rolling and subsequent short term annealing) [13]. Effects of annealing temperature on the crystallographic textures and their correlation with the deformation textures were not reported in that work. The effects of a APCR on the characteristics of oxidation and decarburization of steel have been also reported [14]; results show that columnar grains develop only for rapid decarburization and low oxidation rates. An important issue which was not investigated, is the formation mechanism of the columnar grains during APCR. Therefore, the purpose of the present research is to investigate the microstructural factors that favor this particular grain growth in order to establish their formation mechanism; to this end the effects of annealing prior to cold rolling on phase transformations, grain structure and crystallographic texture were investigated. Changes in crystallographic texture as a function of temperature and their correlation with the deformation textures are also included to show the importance of the development of columnar grains on subsequent processes, and to describe the expected impact on the magnetic properties of steels investigated.

3. Results Since the development of columnar grains occurs only under certain annealing conditions, it is important to investigate the effects of temperature and annealing time on the microstructural changes that take place during thermal treatment, to have a better understanding of their formation mechanism. The following sections describe the effects of annealing prior to cold rolling on phase transformations, grain structure and crystallographic texture of hot-rolled non-oriented electrical steels. Correlation of this later with the deformation textures is also presented to describe the expected impact on the magnetic properties of steels investigated. 3.1. Steel phase transformations Fig. 1 shows the critical transformation temperatures on continuous heating determined by DTA. As can be seen, for the heating rate used in the hot-band annealing (15 °C/min), the critical transformation temperatures, Ac1 and Ac3, are 763 °C and 946 °C, respectively. Therefore, annealing temperatures selected to investigate the microstructural changes within the ferrite (T < Ac1: T < 763 °C), two-phase (Ac1 < T < Ac3: 763 °C < T < 946 °C) and austenite phase field (T˃Ac3: T > 946 °C) were 700–750 °C (α-Fe), 800–900 °C (α-Fe + γFe) and 950–1050 °C (γ-Fe), respectively. Fig. 1 also shows the effect of temperature on the volume fraction of ferrite (α-Fe) and austenite (γFe) calculated by the software JMatPro; as can be observed, the volume fraction of austenite increases while that of ferrite decreases as the temperature is increased above 725 °C (A1). The percentage of these phases at 800 °C is about 87% α-Fe + 13% γ-Fe and varies to 76% αFe + 24% γ-Fe at 850 °C. The proportion of ferrite and austenite at 875 °C is 65% α-Fe + 35% γ-Fe and changes to 50% α-Fe + 50% γ-Fe for a temperature of 900 °C. Finally, the amount of these phases at 950 °C and 1050 °C is about 3% α-Fe + 97% γ-Fe and 0% α-Fe + 100% γ-Fe, respectively. Important to notice is that the temperature calculated by the software to obtain 8% of austenite (γ-Fe8%), is very similar to the Ac1 determined by DTA (770 °C vs 763 °C, respectively), and the temperature calculated by the software to complete the α-Fe to γ-Fe

2. Experimental procedure Table 1 shows the chemical composition of the as-received hotrolled GNO electrical steel band, which was obtained by a local steelmaker. Combustion-infrared absorption spectrometry was used to determine the concentrations of carbon and sulfur according to ASTM E1019, while the content of the other elements was determined by optical emission spectrometry based on ASTM E-403. Chemical composition of the experimental steel was used to calculate both the proportions of phases (ferrite, α-Fe and austenite, γ-Fe) and the equilibrium transition temperatures (A1 and A3). For comparison, the critical transformation temperatures on continuous heating (Ac1 and Ac3) were determined by differential thermal analysis (DTA); to this end samples were heated at 15 °C/min from room temperature up to 1150 °C in an argon atmosphere. Results allow selecting the appropriate annealing temperatures to investigate the effects of annealing treatment on microstructural changes within the single (α-Fe, γ-Fe) and two-phase (αFe + γ-Fe) fields. Hot-rolled bands were pickled by immersion in a 20% HCl solution to remove the oxide produced during hot-rolling process, and then subjected to an annealing treatment. Thermal treatments were carried out in a Thermolyne-6020 muffle-type furnace, using a heating rate of 15 °C/min and air-cooling. Annealing temperatures were 700, 750, 800, 850, 875, 900, 950 and 1050 °C, and soaking times were 10, 30, 60, 90, 120 and 150 min. Microstructure of the as-received and annealed samples was characterized by optical (OM) and scanning electron microscopy (SEM) using an Olympus-XG51 optical microscope and a Philips-XL30 scanning electron microscope, respectively. Average grain size was determined by the linear intercept method as recommended in ASTM E112; fifteen micrographs were analyzed to determine the average value using a commercial image analysis software (Stream Essentials). Crystallographic texture was measured by electron backscatter diffraction (EBSD) and orientation imaging microscopy (OIM). Sample preparation for OIM analysis in normal planes of steel included grinding to half thickness and subsequent polishing to a mirror-like Table 1 Chemical composition of the experimental steel [wt. %]. C

Si

Al

Mn

P

Cr

Cu

S

0.05

0.5

0.21

0.32

0.042

0.019

0.029

0.004

Fig. 1. Evolution of phases as a function of temperature, and transition temperatures determined by JMatPro and DTA. 9

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heating and annealing of steel at temperatures above Ac3, ferrite transforms to austenite and this later can coarsen during heat treatment [16,17]; during cooling, the coarser grains of austenite transform to ferrite with larger size as observed in Fig. 2h and i. Samples annealed at temperatures low within the (α-Fe + γ-Fe) two-phase field region (800 °C and 850 °C), exhibit a particular microstructure consisting on large columnar grains free of carbide particles (Fig. 2d and e). Samples annealed at 875 °C exhibit similar microstructures, but in this case, the microstructure consists on a mixture of columnar plus equiaxed grains (Fig. 2f). Annealing at 900 °C, temperatures high within the two-phase field, leads to equiaxial-grained microstructures (Fig. 2g); the grain size is larger than the one observed in the as-received material or samples annealed at T < Ac1, and it is smaller than the one observed at T > Ac3. Fig. 3 shows the effect of annealing temperature on the average grain size obtained after 150 min of heat treatment. In the case of samples annealed at 875 °C, which exhibited a duplex microstructure (see Fig. 2f), the reported value is the average grain size obtained from measurements of columnar grains. As can be seen, annealing at temperatures below Ac1 causes only a marginal effect on grain size resulting in an average value of about 15 μm (Fig. 3). Annealing at temperatures high within the two-phase field region (e.g. 900 °C) or T > Ac3 (e.g. 1050 °C), results in ferrite grains larger than those observed in the as-received material and in samples annealed at T < Ac1. The maximum grain size obtained at these temperatures is < 30 μm (Fig. 3). In contrast, a significant change in grain size and structure is observed when annealing is carried out at temperatures of 800 °C or 850 °C, in these cases the average grain size is about 625 μm (Fig. 3). As observed in Fig. 4, annealing at 800 °C promotes the development of columnar grains from the surface towards the mid-plane of the sample thickness (Fig. 4); this grain growth occurs isothermally, and thus the size of the columnar grains is larger for a longer time (Fig. 4a, b). Columnar grains grow in a direction parallel to the decarburization direction until they reach the center of steel thickness as observed in Fig. 4c. A similar behavior is observed in samples annealed at 850 °C, but in this case the size of columnar grains is larger for a given time (compare Fig. 4a–c with 4d–f). In the case of samples annealed at 875 °C, the microstructure is in a certain extent similar since columnar grains develop from the surface to the mid-plane thickness. However, for a given time, the size of the columnar grains is even smaller than the one observed at 800 °C or 850 °C (compare Fig. 4g–i with 4a–c and 4d–f, respectively). Annealing at 875 °C for 150 min results in duplex microstructures consisting on large surface columnar grains and small equiaxial grains at the center (Fig. 4i). Fig. 5 shows the average size of columnar grains in samples annealed at 800 °C, 850 °C and 875 °C as a function of time and C concentration. The columnar grain size increases when temperature is increased from 800 °C to 850 °C (Fig. 5a and b); higher temperatures (875 °C) lead to columnar grains with smaller size (Fig. 5c). The rate of growth of columnar grains strongly relates to the decarburization rate (Fig. 5). As can be seen in Fig. 5a and 5b, the size of columnar grains increases significantly with an increase in temperature from 800 °C to 850 °C, this change is accompanied by a considerable reduction in the amount of C. The growth rate of columnar grains in samples annealed at 875 °C is slower than the one observed at 800 °C or 850 °C, which is consistent with the small reduction in the carbon concentration at this temperature (Fig. 5c). Fig. 6 shows the factors that favor the columnar grain growth in hotrolled non-oriented electrical steel bands during annealing prior to cold rolling. Annealing at temperatures within the intercritical range promotes the isothermal γ-Fe to α-Fe transformation due to the steel decarburization (see Fig. 1 and Table 2). Dzubinsky found similar results [18], but during annealing carried out “after cold rolling”. Surface grains with less carbon content growth faster due to lower pinning sites, resulting in larger grains. The advantage in size and lower surface energy of grains resulting from phase transformation allows the preferred

Table 2 Evolution of carbon and oxide characteristics as a function of temperature and time of annealing prior to cold rolling. Characteristics of oxidation/ decarburization process

Time (min) 30 90 150 Oxide thickness Oxide defects

Temperature (°C) Ferrite phase T < Ac1 750

Two-phase field Ac1 < T < Ac3 850

Carbon concentration (wt. 0.050 0.038 0.041 0.013 0.037 0.0092 Thin Thin Porous Cracks

875 %) 0.046 0.032 0.028 Thick Cracked to crack-free

Austenite phase T > Ac3 1050

0.051 0.052 0.053 Thick Crack-free

transformation (A3 = 950 °C) is very similar to the Ac3 determined by DTA (946 °C). Considering the similarity between the theoretical and experimental values of the transformation temperatures, the occurrence of microstructural changes either at temperatures low or at temperatures high within the two-phase field region was determined based on the evolution of phases calculated by the software as a function of temperature. 3.2. Simultaneous decarburization and oxidation The characteristics of the oxidation and decarburization processes that occur during annealing of hot-rolled bands were described in detail in a previous work [14]. However, in order to establish their correlation with the development of columnar grains, some oxide characteristics and carbon concentrations are included in Table 2. Measurements of carbon concentration (hereinafter expressed as weight percent) as a function of temperature and annealing time, indicate that decarburization is slow at temperatures below Ac1. After 150 min of annealing at 850 °C (two-phase field), carbon content is reduced from 0.05% to 0.0092%, which indicates a rapid steel decarburization. Decarburization becomes slower at 875 °C (temperatures high within the two-phase field); in this case, carbon concentration is 0.028% after 150 min. In addition, annealing at temperatures higher than Ac3 makes carbon removal from steel very difficult; as a result, carbon concentration in these samples is very similar to that of the asreceived material (Table 2). Annealing at temperatures below Ac3 promotes the formation of thin, porous or cracked oxide structures, which favor carbon removal. In contrast, thick and crack-free oxide structures obtained after annealing at T > Ac3 make the decarburization process very difficult (Table 2). These results indicate that microchannels such as microcracks or pores favor the carbon removal leading to a rapid steel decarburization [15]. However, their absence make carbon removal from steel very difficult causing a decrease in the decarburization rate as observed in Table 2. 3.3. Microstructure 3.3.1. Size and grain structure Fig. 2 shows the microstructure of hot-rolled strips annealed for 150 min; as can be seen, the size and structure of ferrite grains changes significantly depending on the annealing temperature. Samples annealed at 700 °C and 750 °C (T < Ac1) exhibit equiaxed and fine ferrite grains of around 15 μm; the size and grain structure are very similar to those of the as-received hot-rolled material (Fig. 2a–c), which suggests that annealing at temperatures below Ac1 (763 °C) does not cause a significant effect on microstructure. Annealing at T > Ac3 (946 °C), in the austenite phase field, produces a uniform microstructure consisting on ferrite grains with larger size than those observed in the as-received material or samples annealed at T < Ac1 (Fig. 2h and i). During 10

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Fig. 2. Microstructure of the a) as-received hot-rolled condition, and hot-rolled bands annealed during 150 min at: b) 700 °C, c) 750 °C, d) 800 °C, e) 850 °C, f) 875 °C, g) 900 °C, h) 950 °C and i) 1050 °C.

received hot-rolled material, and in samples annealed at different temperatures during 150 min. Inverse pole figures (IPF) orientation maps use a basic RGB (red, green and blue) coloring scheme, fit to an inverse pole figure [19]. For cubic phases, full red, green, and blue are assigned to grains whose 〈1 0 0〉, 〈1 1 0〉 or 〈1 1 1〉 axes, respectively, are parallel to the projection direction of the IPF [20]. A RGB mixture of the primary components colors intermediate orientations. According to Fig. 7, samples exhibit significant variations in colors meaning differences in the crystallographic texture; this result indicates that annealing temperature can modify this microstructural parameter. Although texture can be discussed from inverse pole figures, the best description of a crystallographic texture in polycrystalline materials can be described quantitatively by the orientation distribution function (ODF) [20]. The ODF is a 3D representation of the “Euler space”, with the three Euler angles (φ1, ϕ, φ2) that describe the orientation of a crystal forming the axes. This function can be represented in the Euler space in three dimensions, however, it is generally interpreted in two dimensions maintaining constant one of the three angles [19]. Fig. 8 shows the φ2 = 45° section of the orientation distribution function (ODF), which presents the ideal orientations in γ-Fe austenite phase that lead to the components of the transformation textures in αFe ferrite phase [21]. The numbers in parentheses correspond to the numbers of variants in γ-Fe phase that can give rise to the same component in α-Fe. As can be seen in this Figure, more than one orientation (variant) in γ-Fe before phase transformation can lead to the same orientation in α-Fe. It is for this reason that it is difficult to determine with accuracy the origin of the transformation textures [21]. The Goss (110)[001], rotated Goss (110)[11¯0] and rotated cube (001)[11¯0], (001)[11¯0] texture components in ferrite form from the cube (001)[11¯0] in the parent phase (austenite). In addition, the rotated cube

Fig. 3. Effect of temperature on the average grain size of hot-rolled bands after 150 min of annealing.

growth in a direction parallel to the decarburization direction (Fig. 6). Kovac reported similar results [8] during decarburizing annealing of non-oriented electrical steels carried out “after cold rolling”.

3.3.2. Crystallographic textures produced by annealing prior to cold rolling and their correlation with the deformation textures Crystallographic texture also changes significantly with annealing temperature. Fig. 7 shows the orientation maps obtained in the as 11

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Fig. 4. Evolution of columnar grain growth during isothermal annealing: a, b, c) 800 °C, d, e, f) 850 °C and g, h, i) 875 °C, during: left) 30 min, center) 90 min and right) 150 min.

components also form from the Br (110)[11¯2] component [21]. During the γ-Fe → α-Fe transformation, the Cu component is replaced by the transformed Cu. Additionally, the Br component is replaced by the transformed Br (554)[225¯] to (332)[1¯1¯3], rotated cube (001)[11¯0] and (001)[1¯1¯0] components, and an additional orientation located in the proximity of the (112)[13¯1] component. Other components in austenite such as the S (123)[634¯], Goss/Br (011)[51¯1], Br/S (168)[211¯] and S/Cu (236)[322¯], can also contribute significantly with the transformation texture in the α-Fe phase [21]. Fig. 9 shows the φ2 = 45° section of the ODF corresponding to the as-received hot-rolled (Fig. 9a) and annealed samples (Fig. 9b, c, d). The main texture components observed in the as-received material are: (001)[1¯1¯0], (001)[11¯0], (111)[11¯0] and (111)[01¯1] (Fig. 9a). Annealing at temperatures within the ferrite phase field (700 °C) results in a random texture characterized by four components: (001)[11¯0], (001)[1¯1¯0], (223)[11¯0] and (111)[01¯1¯] (Fig. 9b). Annealing at temperatures within the austenite phase field (1050 °C), causes an increase in the intensity of the (001)[11¯0] and (001)[1¯1¯0] components (Fig. 9d) and development of the (112)[11¯0], (111)[11¯0], (111)[01¯1] and (111)[1¯1¯2] components (Fig. 9d). The most important observation regarding the crystallographic textures obtained at temperatures within the two-phase field (850 °C) is the development of a set of orientations near to the cube texture (001)[01¯0], and a spread with a maximum intensity in the vicinity of the (114)[48¯1] (deviation about 20° from (113)[11¯0]). The following components are also distinguished: (331)[1¯1¯6] and (116)[3¯3¯1] (component deviated 12.4° from (001)[1¯1¯0]) . In addition, it can be seen that annealing at this temperature does not cause the development of components along the γ-fibre, 〈1 1 1〉//ND (normal direction) (Fig. 9c). Fig. 10 shows a comparison between the deformation textures of

samples processed by the conventional route (without annealing, Fig. 10a) and samples previously annealed during 150 min at 700 °C (Fig. 10b), 850 °C (Fig. 10c) and 1050 °C (Fig. 10d), temperatures within the ferrite phase field, two-phase region, and austenite phase field, respectively. In general, the deformation texture of samples without previous annealing and samples previously annealed at 700 °C exhibits an increase in the intensities of both the α-fibre (〈1 1 0〉//RD, RD = rolling direction) and γ-fibre (〈1 1 1〉//ND, ND = normal direction). Deformation textures also show an increase in the intensity of (001)[1¯1¯0], and the absence of the (001)[11¯0] component both observed in the as-received hot-rolled material (compare Fig. 9a with Fig. 10a and b). The texture components with the highest intensity in the cold-rolled condition of samples without previous annealing were (223)[11¯0] and (111)[1¯1¯2] (Fig. 10a). The main texture components developed in samples annealed at 700 °C were (111)[1¯1¯2] and a spread from (223)[11¯0] to (111)[12¯1] (Fig. 10b). In contrast, the deformation texture of samples previously annealed at 850 °C mainly exhibits orientations along the θfibre, 〈0 0 1〉//ND. In this case, the texture components with highest intensity are (001)[11¯0], (001)[01¯0] and (001)[1¯1¯0]. In addition, there is a significant decrease in the intensity of the γ-fibre (Fig. 10c). Finally, annealing at 1050 °C causes a significant increase in the intensity of the following components: (111)[11¯0], (111)[01¯1], (112)[11¯0], (114)[11¯0], (001)[5¯7¯0] (001)[1¯1¯0]) (deviation of 10.2° from and (001)[23¯0] (deviation of 11.2° from (001)[11¯0]) (Fig. 10d).

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(Fig. 4). Normal grain growth results from the interaction between topological requirements of space-filling and the geometrical needs of surface tension equilibrium. Factors affecting this type of grain growth are energy at grain boundaries, radius of curvature and boundaries mobility. Abnormal grain growth originates by the preferential growth of a few grains, which have some special growth advantage over their neighbours; this grain growth is strongly affected by second-phase particles, texture and surface effects [22]. Abnormal grain growth is particularly favored if the texture leads to a significant variation of surface energy among the grains. This type of growth is possible in such a situation because within a highly textured volume the grain boundaries have a lower misorientation and hence a lower energy and mobility than within a normal grain structure. The occurrence of abnormal grain growth may be limited by “nucleation” rather than growth considerations. Strong evidence for this comes from the large number of experimental observations which show that abnormal grain growth is particularly likely to occur as the annealing temperature is raised and as the particle dispersion becomes unstable [22,23]. Coarsening of the particle dispersion or a lowering of the volume fraction should not in itself make abnormal grain growth more likely. Therefore, what probably occurs is local destabilization of the grain structure by removal or weakening of critical point, leading to a local inhomogeneous grain growth. The resulting broader grain size distribution then enables the “nucleation” of abnormal grain growth. As the driving force for grain growth is usually very small, significant grain growth is often found at very high temperatures [22,23]. The results of the present work show that abnormal grain growth is not presented necessarily at the highest temperatures as observed in Figs. 2 and 4; this particular grain growth occurs only at temperatures low within the twophase field region. The effect of a preferred texture on the development of columnar grains is also discarded considering that the as-received material exhibits a random texture (Fig. 9a), and that textures resulting from annealing at 850 °C are completely different from those observed in the as-received condition. In conclusion, results show that the development of columnar grains in non-oriented electrical steels subjected to annealing prior to cold rolling occurs by a process of chemical destabilization of the two-phase microstructure (Fig. 1) due to decarburization of steel under isothermal conditions (Fig. 4). During the ferrite to austenite (α-Fe → γ-Fe) transformation, dissolution of carbides occurs [24]. In this process, the carbon atoms migrate to the austenite through the γ-Fe/α-Fe interface due to the lower carbon solubility in ferrite and then diffuse into the austenite grains forming a solid solution. Therefore, the few austenite formed when annealing is carried out at temperatures low within the two-phase field region (α-Fe + γ-Fe) is enriched in carbon [24]. Decarburization of γ-Fe grains formed during the intercritical annealing (Fig. 1) promotes the isothermal γ-Fe → α-Fe phase transformation, which may cause a local destabilization of the grain structure by removal of carbides leading to a local inhomogeneous grain growth at the steel surface. Carbides removal will reduce the pinning points leading to surface ferrite grains with larger size and lower carbon content (Table 2, Fig. 6), which grow preferentially against the C concentration gradient (Fig. 4). With increments in the annealing time, the decarburization progresses as a front moving from the surface towards the two-phase middle region of the strip (Fig. 4). The formation of the thin surface layer of ferrite grains due to the rapid steel decarburization produces a discontinuous concentration profile of carbon through the steel thickness. The carbon content is low at the surface (decarburized grains) and high at the α-Fe/(α-Fe + γ-Fe) interface, and even below the interface due to the presence of carbides (Fig. 6). These results are consistent with the development of columnar grains during a stepped temperature annealing [8]. The remaining austenite grains in the central region of the strip, which are rich in carbon and prevent grain growth, also transform to ferrite with the progress of decarburization until the fine and equiaxial

Fig. 5. Evolution of columnar grain size and carbon concentration during annealing at: a) 800 °C, b) 850 °C and c) 875 °C.

4. Discussion 4.1. Columnar grain growth assisted by decarburization and phase transformation Grain growth may be divided into two types, normal grain growth and abnormal grain growth. During normal grain growth, the microstructure changes in a rather uniform way, there is a relatively narrow range of grain sizes and shapes, and the form of the grain size distribution is usually independent of time [22]. There are however, circumstances when the microstructure becomes unstable and a few grains grow excessively, consuming the matrix of smaller grains and resulting in a bimodal grain size distribution; this process is known as abnormal grain growth [22]. According to the microstructures presented in Fig. 2, annealing at 700, 750, 900, 950 and 1050 °C causes a uniform grain growth maintaining the equiaxial shape observed in the as-received material, which are some characteristics of normal grain growth. In contrast, annealing at 800, 850 and 875 °C promotes an abnormal grain growth characterized by a discontinuous grain growth and the formation of duplex microstructures; in this case, columnar grains consume the matrix of smaller grains leading to a bimodal grain size distribution

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Fig. 6. Microstructural factors that favor the columnar grain growth in non-oriented electrical steels.

determined the diffusion coefficients for the ferrite and the austenite during decarburization of electrical steels. Evaluation of these coefficients at 900 °C results in a value of 39 times smaller for the γ-Fe phase. Therefore, the diffusion process in this phase would be slowed in the same proportion [26]. The austenite in the (α-Fe + γ-Fe) field acts as a pinning point and

matrix is consumed (Figs. 4 and 6). The presence mainly of ferrite in the two-phase (α-Fe + γ-Fe) microstructure (Figs. 1 and 6) favors the columnar grain growth during the simultaneous oxidation/decarburization process. Activation energy of carbon atom diffusion in austenite is significantly higher than in ferrite, which indicates that carbon diffusion rate in ferrite is faster than in austenite [25]. Carlos [26],

Fig. 7. Orientation maps corresponding to: a) as-received hot-rolled samples, and samples annealed for 150 min at: b) 700 °C, c) 850 °C and d) 1050 °C. 14

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Fig. 8. φ2 = 45° section of Euler space showing common texture components and fibres formed during rolling and annealing of BCC metals and alloys. The α-Fe texture components formed from the γ-Fe components are included [21].

Fig. 9. φ2 = 45° section of the orientation distribution function: a) as-received hot-rolled samples, and samples annealed for 150 min at: b) 700 °C, c) 850 °C and d) 1050 °C. 15

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Fig. 10. Deformation textures obtained after cold-rolling: a) samples without annealing, and samples previously annealed for 150 min at: b) 700 °C, c) 850 °C and d) 1050 °C.

very difficult reducing the driving force for this particular grain growth [27]. The absence of columnar grains at T ≥ 900 °C (temperatures high within the two phase field region), could be attributed to the combination of both the higher oxidation rate and the higher amount of austenite (higher carbon solubility) in the two-phase field region. These two factors slow down the rate of decarburization and minimize the driving force for this particular columnar grain growth [28].

prevents growth but the most deleterious effect is the decarburization delay because of the austenite very low diffusion coefficient. Anchors removal is due to the decarburization process where austenite transforms into ferrite [26]. According to this, it is expected that columnar grain growth in the intercritical region is faster for lower volume fractions of γ-Fe as observed in samples annealed at 800 °C and 850 °C (see Figs. 1–5 and Table 2). As the amount of austenite increases, carbon diffusion decreases due to its higher solubility leading to a reduction of the rate at which columnar grains grow as observed in samples annealed at 875 °C (see Figs. 1, 2, 5 and Table 2). Therefore, the concentration gradient generated in the microstructure due to the different solubility of carbon in ferrite (α-Fe) or austenite (γ-Fe), which is dependent on the proportion of phases, represents the driving force for the columnar grain growth.

4.3. Evolution of crystallographic texture The crystallographic texture of samples annealed at 700 °C (Fig. 9b) and 1050 °C (Fig. 9d) is very similar to that of the received material (Fig. 9a). Annealing at temperatures within the single ferrite (below Ac1) or austenite (above Ac3) phase fields does not have a significant effect on texture. The main texture components observed in these samples were (001)[1¯1¯0], (001)[11¯0] and a spread along the γ-fibre, (111) //NP (Fig. 9a, b and d). Changes in texture of hot-rolled bands produced by annealed at 700 °C and 1050 °C can be explained in terms of the oriented nucleation mechanism, in which growing grains develop crystallographic textures similar to those of grains from which they grow [29]. Annealing at temperatures within the intercritical region (850 °C) promotes the development of columnar microstructures and causes a significant change in the crystallographic texture. For instance, components (111)[11¯0] and (111)[01¯1] observed in the as-received material (Fig. 9a) are not present after annealing at 850 °C (Fig. 9c). In addition, the component (001)[01¯0] which was not present in the asreceived material, is observed after annealing at 850 °C (compare Fig. 9a and c). Ray [30], mentioned that the final texture obtained after hot-rolling

4.2. Influence of oxidation on the columnar grain growth Formation of the columnar microstructures would be expected when the velocity of the transformation front is higher than the nucleation speed of the phase transformation [8]. Oxidation of steel affects the decarburization rate and thus the kinetic of the isothermal (γFe → α-Fe) austenite to ferrite phase transformation (Fig. 4 and Table 2); as a result, it also has a negative effect on the rate at which columnar grains grow, as observed in samples annealed at 875 °C (Fig. 5). The absence of columnar grains at temperatures within the single ferrite or austenite phase fields relates to the low decarburization rate, influenced by lower carbon diffusion at T < Ac1 and higher oxidation rates at T > Ac3. The two conditions make carbon removal from steel 16

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of steel in the two-phase field region (α-Fe + γ-Fe) must be that resulting from the γ-Fe → α-Fe phase transformation, and other fraction related to the presence of ferrite phase. In this context, these results suggest that the nucleation phenomenon in this case is clearly dominated by a selective growth that takes place during the austenite to ferrite transformation, since neither of the textures present in samples annealed at 850 °C could be reproduced solely on the basis of oriented nucleation [29]. Although changes in texture can be explained in terms of the Bain ({0 0 1}γ||{0 0 1}α : 〈1 0 0〉 γ||〈1 1 0〉 α) and KurdjumovSachs (K-S) ({1 1 1}γ||{1 1 0}α : 〈0 1 1〉 γ||〈1 1 1〉 α) correspondence relationships [21], it is clear that the rate of decarburization also plays an important role on both grain structure and resulting texture. Textures by “oriented nucleation” develop when normal grain growth takes place (Fig. 9b and d), it means for low decarburization rates. In contrast, rapid decarburization promotes a selective growth with a significant change in texture (Fig. 9c). Li [31] reported that iron oxidation or oxidation of alloying elements on the surface hinders the decarburization process and, consequently, the preferred microstructure and texture evolution in non-oriented electrical steels. The results of the present work suggest that decarburization, which has a strong dependence on the oxidation rate, affects strongly the texture obtained after annealing prior to cold rolling, which is in agreement with the results obtained by Li [31]. On the other hand, it has been reported that the rolling texture of iron and low carbon steels is largely independent of composition; in general, five orientations describe the components of this texture: {111}112 , {111}123, {001}110 , {112}110 and {111}110 [22,23]. The usual description has concentrated on two prominent orientation spreads, which are in turn, described as fibre textures. One of these corresponds to a 111 fibre axis perpendicular to the sheet surface; the above {111} uvw orientations are prominent in this spread, which is well known as γfibre [22,23] (Fig. 8). The other is a partial fibre texture with a 110 fibre axis parallel to the rolling direction; the {hkl}110 orientations are prominent in this spread, which is known as α-fibre [22,23] (Fig. 8). In contrast the θ-fibre (Fig. 8), {001} uvw , usually decreases during rolling of iron and low carbon steel. As observed in Fig. 10, the deformation textures of samples without previous annealing (Fig. 10a) and samples previously annealed at 700 °C (Fig. 10b) and 1050 °C (Fig. 10d), exhibit an increase in the intensity of the α and γ fibres as well as a decrease in the intensity of the (001)[1¯1¯0] component, compared with intensities observed in the asreceived hot-rolled condition (Fig. 9a). Some of the components mentioned above are observed in the ODF’s of deformed samples (Fig. 10a, 10b and 10d). The deformation texture of samples previously annealed at 1050 °C (T > Ac3) show a significant increase in the intensity of the (111)[11¯0], (111)[01¯1], (111)[1¯1¯2] and two components near the (001)[11¯0] and (001)[1¯1¯0] (Fig. 10d), which relates to the increase in the intensity of these componentes due to the γ-Fe to α-Fe phase transformation (Fig. 9d). The deformation texture of samples previously annealed at 850 °C is completely different; in this case, the texture is characterized by the absence of α and γ fibres and the development of a strong θ-fibre (Fig. 10c). These effects relate to the development of columnar-grained microstructures during annealing prior to cold rolling. As observed in Fig. 9c the development of columnar grains results in a considerable reduction in the intensity of the α-fibre and the γ-fibre, an increase in the intensity of (001)[01¯0] and a component near the (001)[1¯1¯0], and a spread in the vicinity of the θ-fibre. Kestens et al. [32] reported that rotation paths of the orientation flow during plastic deformation is affected by changing the orientation of the rolling-mill rolls with respect to the steel sheet. In this way, the conventional deformation texture of non-oriented electrical steels obtained by the typical process is modified to produce an extremely strong rotated cube component {001}110 with an extraordinary intensity, without the development of α and γ fibres. In the study, the hot-rolled steel strip was rotated 90° with respect to the rolling direction before

cold rolling. In doing so the rolling direction of hot-rolling turns into the transverse direction of cold rolling and vice versa [32]. Hence, the hot-band texture with its characteristic α and γ fibres is transformed into a texture with a strong θ-fibre by simply substituting the hot rolling by the cold rolling reference frame. Further annealing of cold-rolled samples lead to recrystallization textures very similar to the ones developed during cold-rolling, which were characterized by a spread in the vicinity of the θ-fibre [32]. The results obtained in the present investigation show that annealing prior to cold rolling represents an alternative method to modify the conventional deformation texture of non-oriented electrical steels obtained by the typical process. This is related to variations in rotation paths of the orientation flow during plastic deformation, due to changes in the crystallographic texture produced by the development of columnar grains during annealing prior to cold rolling. At room temperature, iron has a body centered cubic lattice and within a crystal of iron in which electron spins have spontaneously aligned to give self-saturation, some directions are much easily magnetisable than others. In a BCC lattice, [111] is a body diagonal, [110] is a face diagonal and [100] is a cube edge; the [100], [110] and [111] are the easy, intermediate and hard directions of magnetisation [33]. A positive impact in the magnetic properties of the final product is then expected considering the presence of θ-fibre (001//ND) and the absence γ-fibre (111//ND) in the cold-rolled in the cold-rolled condition, which according to Kestens et al. [32], could lead to the development of favorable texture components during subsequent annealing treatment enhancing the magnetic behavior of the investigated steels. 5. Conclusions The following conclusions can be drawn from the results obtained in the present work:

• The development of columnar grains by annealing prior to cold •

• •

•

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rolling is an efficient method to modify grain structure and crystallographic texture of hot-rolled non-oriented electrical steels. Decarburization of γ-Fe grains formed during intercritical annealing promotes the isothermal γ-Fe → α-Fe phase transformation, which may cause a local destabilization of the grain structure by removal of carbides leading to a local inhomogeneous grain growth at steel surface. Carbides removal will reduce the pinning points favoring boundary motion and leading to surface ferrite grains with larger size and lower carbon content. The advantage in size and lower surface energy of grains resulting from phase transformation allows the preferred growth in a direction parallel to the decarburization direction. The Bain and Kurdjumov-Sachs correspondence relationships explain the changes in the crystallographic texture. Oxidation of steel affects the decarburization rate and have a strong effect on the resulting texture. Annealing prior to cold rolling represents an alternative method to modify the conventional deformation texture of non-oriented electrical steels obtained by the typical process. A positive impact in the magnetic properties of the final product could be expected considering the presence of θ-fibre and the absence γ-fibre in the coldrolled condition, which could lead to the development of favorable texture componentes during the subsequent annealing treatment, enhancing the magnetic behavior of the investigated steels. Decarburization annealing in the conventional processing route is carried out after cold rolling at T < Ac1 during times longer than 16 h. Therefore, annealing prior to cold rolling at temperatures within the two phase field region (Ac1 < T < Ac3) during 150 min, could be an attractive method not only to obtain the microstructural characteristics required in these steels prior to cold rolling, but also to increase productivity of non-oriented electrical steels.

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Acknowledgements

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E. Gutiérrez Castañeda recognizes the Cátedra CONACYT assigned at the Institute of Metallurgy of the Autonomous University of San Luis Potosi (IM-UASLP). Authors of the present research would like to thank all the facilities given at CINVESTAV-IPN and at the IM-UASLP to carry out the experimental work. The valuable time and technical assistance of Teodoro Caballero, Rosa Lina Tovar Tovar and Nubia Arteaga Larios is duly recognized. First author would like to thank the “Comisión de Investigación y Desarrollo Tecnológico (CIDT)” and “Secretaría de Investigación y Posgrado (SIP)” of the Autonomous University of San Luis Potosi, for the Fund for Research Support FAI-UASLP 2016 through the project C16-FAI-09-62.62. Data Availability Data will be made available on request. References [1] C. Nepthali, J.S. Ma, S. Armando, J.G. Emmanuel, A.R. Iván, R.C. Francisco, Influence of thickness and chemical composition of hot-rolled bands on the final microstructure and magnetic properties of non-oriented electrical steel sheets subjected to two different decarburizing atmospheres, Metals 229 (2017) 1–16. [2] J.M. Anthony, Energy efficient electrical steels: magnetic performance prediction and optimization, Scr. Mater. 67 (2012) 560–565. [3] S. Jabobs, D. Hectors, F. Henrotte, M. Hanfer, M. Herranz, K. Hameyer, P. Goes, Magnetic material optimization for hybrid vehicle PMSM drive, World Electr. Veh. J. 3 (2009) 1–9. [4] A. Chaudry, R. Khatirkar, N.N. Viswanathan, V. Singal, A. Ingle, S. Joshi, I. Samajdar, Low silicon non-grain-oriented electrical steel: linking magnetic properties with metallurgical factors, J. Magn. Magn. Mater. 313 (2007) 21–28. [5] B. Philip, Electrical Steels for Rotating Machines, Manufacturing Methods, The institution of Electrical Engineers, IEE, London, 2002. [6] K. Verbeken, E. Gomes, J. Schneider, Y. Houbaert, Correlation between the magnetic properties and the crystallographic texture during the processing of non-oriented electrical steel, J. Solid State Phenom. 160 (2010) 189–194. [7] Darja S., 2010. Non-oriented electrical steel sheets. MTAEC9. 44, pp. 317–325. [8] F. Kovac, M. Dzubinsky, Y. Sidor, Columnar grain growth in non-oriented electrical steel, J. Magn. Magn. Mater. 269 (2004) 333–340. [9] Y. Sidor, F. Kovac, V. Petrychka, Secondary recrystallization in non-oriented electrical steels, METABK 44 (2005) 169–174. [10] I. Petryshynets, F. Kovac, M. Molnarova, P. Gavendova, M. Sopko, B. Petrov, Columnar grain growth with enhanced rotation texture in temper rolled NO silicon steels, Mater. Sci. Forum. 782 (2014) 201–204.

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