Two Mo-Nb microalloyed transformation-induced-plasticity (TRIP) steels, with Al contents of 0.23 and 0.65, were subjected to several hot-rolling conditions designed to generate different ferrite morphologies and grain sizes. These structures were then cold rolled and TRIP annealed under different heat-treatment conditions. To further develop TRIP steel in terms of strength and ductility, stabilizing retained austenite by isothermal bainitic transformation was studied in detail. Microstructure observation and tensile tests were conducted, and volume fractions of retained austenite were measured. It was observed that increasing the aluminum content enhances the transformation rate and increases the total amount of bainite fraction at the expense of retained austenite. The latter effect enhances formability by increasing ductility. Furthermore, it was observed that the hot-rolling schedule, prior to cold rolling and heat treatment, has a decisive effect on structure refinement, which enhances the strength-ductility balance of the final product. To study the transformation behavior, dilatometer testing was conducted under conditions similar to that of the heat treatment. Thermodynamic calculations were used to verify the results. DOI: 10.1007/s11661-008-9594-2

(c) The Minerals, Metals & Materials Society and ASM International 2008

I. INTRODUCTION

THE remarkable strength-ductility balance in transformation- induced-plasticity (TRIP) steel results from the occurrence of the TRIP phenomenon during deformation[1] The coexistence of austenite with a certain microstructural stability is of vital importance in order for this phenomenon to occur and, hence, to achieve the desired properties. The austenite retention is usually obtained by combining the effects of chemical composition and typical heat treatment. In this respect, adding large amounts of silicon to TRIP steel ensures that cementite precipitation is unlikely to occur in the microstructure during bainite formation.[2) The absence of cementite ensures that the carbon will enrich the austenite rather than form cementite plates. Therefore, after the bainite transformation finishes by further cooling to room temperature (RT), the austenite is stabilized. Jeong et al. concluded that austenite retention in these low-alloyed steels is almost impossible with silicon concentrations much below 1 wt pct.[3] However, these high required silicon levels are outside standard industrial practice for producing flat products because of the following.

(a) Steel with more than 1 pct Si has a poor Zn coating quality after continuous galvanizing, due to the presence of Si-Mn oxides on the strip surface.[4]

(b) The high Si content of these steels causes red scales to form in bands. After pickling, the oxides are completely removed, but the band remains visible on the surface of the pickled steel.[5]

Consequently, studies have been performed with other elements that can substitute for the role of silicon. Presently, aluminum seems to be the most promising candidate. However, in addition to the fact that high aluminum content in steel causes serious casting problems, a full substitution of silicon by an equivalent amount of aluminum leads to a marked deterioration of the strength-ductility balance.[6,7] De-Meyer et al.[8] proposed that silicon can best be partially replaced by aluminum with an increase in the carbon content. Bleck also suggested that combining silicon, aluminum, and phosphorus is a reasonable compromise and could be the most important alloying concept for low-alloyed TRIP steels.[9]

Little or no information can be found in the literature about the sole effect of aluminum content on the phase transformation and mechanical properties of the TRIP-aided steels. De-Meyer et al.[8] and E. Emadoddin et al.[10] presented some results about Si-Al- alloyed TRIP steels. However, their comparison of the number of phases and the mechanical properties is misleading, because the carbon content was different in their alloys. Carbon is the main alloying element by which both all transformations are noticeably affected and the final microstructure and the mechanical properties are controlled.[11] On the other hand, without changing the carbon content of the alloys, Jacques et al.[7] studied the effect of a partial substitution of Si by Al while varying the silicon content.

The present study is aimed at ascertaining how Al content variations, the prior hot-rolling conditions, and heat-treatment parameters affect the microstructure and mechanical properties of the cold-rolled Mn-Si-Al TRIP-aided steel alloyed with Mo-Nb.

II. EXPERIMENTAL PROCEDURE

A. investigated Materials

The Si-Al-Mo-Nb steel studied in this work was produced in the laboratory. The two alloys studied, steels 1 and 2, differ in their Al content. The chemical composition is given in Table I. The alloys contain Mo and are microalloyed with Nb. The Nb in solid solution has been found to improve the TRIP properties; the Mo retards the precipitation of Nb(C,N), thus potentially improving the effectiveness of Nb as a TRIP enhancer [12.13] furthermore, Mo retards austenite transformation to both ferrite and pearlite, effecting more manageable process control.[14]

B. Heat Treatment and Dilatometry

The heat treatment had been conducted on the mechanical testing samples and on samples for microconstituents investigations using salt baths. This had been done by austenitizing in Durferrit GS 540/ R2 and austempering in Durferrit AS 140 salt baths*. The inertor R2 was added, to prevent any oxidation or decarburization during austenitizing.

* Durferril GS 540/R2 and Durferrit AS 140 are trademarks of Durferrit GmbH, Mannheim, Germany.

Dilatometric measurements were conducted on a Baehr dilatometer DIL 805A/D** (Figure 1), which has a resolution of 0.05 [mu]m/0.05 K. All the dilatometric measurements were performed using specimens 2.5 x 5 x 10 mm in size. The test specimens were degreased with an acetone solvent. Sheathed type S Pt/Pt-10 pct Rh thermocouples with a nominal diameter of 0.1 mm were individually spot welded to the surface of the specimen in the central position of the 5 x 10- mm surface-to-monitor temperature. Each sample was held between two quartz rods, with its 10-mm side along the rods. One of the rods is fixed; the other one is connected to a linear variable differential transducer (LVDT). A reference rod is also connected to the LVDT. The dimension variations of the specimens during the thermal cycle are transmitted via the moving quartz pushrod to the LVDT sensor. After placing the sample between the pushrods, the insulating sheaths on the thermocouple wires had been moved along the thermocouple wires until they contacted the specimen surface.[15] The thermal cycles were performed under a vacuum of 0.005 Pa, and helium was used for cooling. A computer and data acquisition system recorded the dilatometric change and temperature as a function of time, and cross correlated the relative change in length as a function of temperature.

** Baehr Dilatometer DIL 805A/D is a trademark of Bahr- Thermoanalyse GmbH, Huellhorst, Germany.

C. Materials Characterization

The materials were machined to the required specimen size and geometry prior to heat treatment. This was done to avoid the transformation of retained austenite to martensite due to machining forces. Standard subsize tensile specimens were machined in accordance with ASTM standard E8-03; the specimens had a width of 6.4 mm and a gage length of 25.4 mm and were machined transverse to the rolling direction from the cold-rolled bands.

For investigating micro structural constituents, the samples were ground and polished using the normal metallographic preparation procedure. The microstructures were examined with an optical microscope after etching, using LePera[16,17] or nital etchant. The samples for the scanning electron microscope (SEM) were tempered at 473 K for 2 hours before mechanical preparation and deep etching in nital. The tempering treatment at 2 hours and 473 K was performed to enable a good resolution of the martensite substructure.[16]

A magnetic measurement technique that made use of a hysteresis recorder was employed to estimate the amount of the retained austenite (K^sub gamma^). As compared to X-ray diffraction (XRD), the advantages of the magnetic technique are that magnetic measurement is performed on the whole volume, the specimens require no special preparations, and the measurements are fast, reliable, and more sensitive to retained austenite.[18,19]

The thermodynamic calculations were performed using THERMO- CALC[dagger] and the database TCFE3; the ferrite, austenite, and cementite phases were considered.

[dagger] THERMO-CALC is a trademark of Thermo-Cale software AB, Stockholm, Sweden.

III. RESULTS AND DISCUSSION

A. Hot-Rolling Conditions

1. Estimation of nonrecrystallization temperature

When dynamic recrystallization (DRX) takes place during material deformation, e.g., a rolling process, grain size is determined by the steady-state flow stress. Whenever the critical strain epsilon^sub c^ for the onset of DRX is reached and exceeded during hot deformation, a metadynamic recrystallization (MDRX) takes place after the straining is interrupted, consequently coarsening the austenite grains. If deformation is interrupted before reaching the strain epsilon^sub c^ and if the temperature is high enough, then static recrystallization (SRX) takes place. The austenite grains are undergoing refinement by deformation in the recrystallization region. This is important because the grain size of the austenite strongly affects both the kinetics of the subsequent gamma [arrow right] alpha transformation as well as the ferrite grain size, i.e., smaller austenite grains lead to smaller ferrite grains. When deformations are applied at temperatures below that of nonrecrystallization (T^sub nRX^), the austenite grains elongate and deformation bands form within the grains. The process is called "pancaking." As the deformation amount increases in this region, the number of nucleation sites at the austenite grain boundaries and within the austenite grains increases. Thus, the gamma [arrow right] alpha transformation from deformed austenite yields much finer ferrite grains than does the transformation from recrystallized and strain-free austenite. Therefore, T^sub nRX^ is a very important parameter and its determination represents a crucial step in designing rolling schedules. Accordingly, the present hot-rolling schedules were selected according to this temperature. The T^sub nRX^ was determined using the method described elsewhere.[20-22] The estimated values of T^sub nRX^ for steels 1 and 2 are 1138 and 1159 K, respectively. 2. Hot-rolling schedules

With these estimated temperatures, the deformation part for the schedules was determined in such a way that all three possibilities were covered, namely, all deformations conducted above T^sub nRX^, deformations conducted below T^sub nRX^, and deformations conducted at mixtures of temperatures below and above T^sub nRX^. Figure 2 shows a schematic representation of the employed schedules. The schedule that resulted in microstructure formed from the recrystallized austenite is denoted R, whereas that which resulted in microstructure formed from pancaked austenite is denoted P. The RP schedule is for the microstructure resulting from the recrystallized and then pancaked austenite.

A reheating temperature of 1523 K was selected to dissolve the Nb in solution entirely, as shown in Figure 3. In this figure, the Nb in austenite increases by increasing the reheating temperature, so that it finally reaches its bulk content in the alloy at approximately 1523 K. If the homogeneous annealing temperature prior to hot rolling is less than the dissolving temperature of Nb(C,N), the undissolved carbonitrides will exist and abnormally grow so that they weaken the refinement of austenite grains and the precipitation strengthening of the Nb(C,N).[23]

Thus, during reheating before hot rolling, the Nb completely dissolves in austenite; then, during the subsequent cooling and hot- rolling process, the Nb(C1N) precipitates. The addition of Nb to TRIP steels effectively refines the austenite grain during the hot- rolling process because the precipitates retard austenite recrystallization and, in turn, refine the final microstructure after intercritical annealing. The recrystallization process is prevented by copious precipitation below T^sub nRX^.[24,25] An extensive study of the precipitation processes in Nb-microalloyed TRIP-steel is reported elsewhere.[23]

The cast ingots were hot rolled in four passes (Figure 2), from a thickness of 19 to 4 mm, and with a true strain value phi = 0.38 at each pass. The temperature was continually monitored with a pyrometer during the final cooling of the hot-rolled strips in air. It took between 20 and 31 minutes to cool the strips from the finish rolling temperature to 623 K. The surface oxide scale was then removed from the slabs using shot blasting; finally the 4-mm-thick hot-rolled plates were cold rolled to a thickness of 2.5 mm.

Figure 4 shows the microstructure obtained after the different hot-rolling schedules. Acicular ferrite (AF) dominates and polygonal ferrite (PF) is also present. Vickers microhardness tests HV0.05 were performed within regions occupied exclusively by PF and AF (Figures 5(a) and (b)). The average HV0.05 values for PF and AF phases are 1370 and 2900 MPa, respectively.

The AF is a nonequiaxed ferrite (Figure 5(c)) formed upon continuous cooling by a mixed diffusion and shear mode of transformation that begins at a temperature slightly higher than the transformation temperature of upper bainite. The Mo and Mn elements (Table I) causes the beginning of transformation to be retarded, because the ferrite start curve on the continuous cooling transformation (CCT) diagram is displaced to longer times. Therefore, the austenite decomposes at lower temperatures during cooling, yielding fine-grained, AF.[26,27]

Figure 5(c) shows a remarkable characteristic of this type of microstructure: it possesses an irregular configuration, which has various grain sizes distributed in a heterogeneous manner with random orientations. The PF grains are equiaxed with different average sizes, depending on the applied rolling schedule.

Table II compares the PF contents (ferrite percent) and their grain sizes (d^sub f^) that resulted from the different hot-rolling schedules for both alloys. The higher ferrite content of steel 2 can be explained by its lower hardenability, because its higher aluminum content shifts the ferrite nose in the CCT diagram toward a lower time.[11] It is also observed that the hardenability is sensitive to the deformation temperature. Decreasing the deformation temperature decreases the steel hardenability; hence, the ferrite content increases by moving the deformation schedule toward a lower temperature.

On the other hand, the variation in the ferrite grain size that occurs when the hot-rolling schedule is varied is caused by the different austenite grain morphology that forms the ferrite. For the P schedule, the grains are produced from work-hardened (pancaked) austenite through an austenite-to-ferrite transformation. The sizes of these ferrite grains are smaller than those produced from the recovered and recrystallized austenite grains of the R schedule. It is well known that ferrite tends to nucleate at the austenite grain boundaries.[28'29] Thus, the ferrite grain size developed after the transformation strongly depends upon the austenite grain structure that appears just before the start of the transformation. During the straining of the austenite in the nonrecrystallization region, deformation bands and twinning boundaries form and the dislocation density inside austenite grains is greatly increased; the increased dislocation density provided favorable nucleation sites and an enhanced nucleation rate.