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The hardenability of austenite with different alloy content and dispersion in dual-phase steels
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 8 ( 2 0 0 8 ) 75–84
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The hardenability of austenite with different alloy content and dispersion in dual-phase steels ˘ b,∗ B. Demir a , M. Erdogan a
Materials Division, Metallurgy Education Department, Faculty of Technical Education, Karaelmas University, 78050 Karabuk, Turkey b Materials Division, Metallurgy Education Department, Faculty of Technical Education, Gazi University, 06500 Besevler, Ankara, Turkey
a r t i c l e
i n f o
a b s t r a c t
Article history:
In this investigation, two steels with different chemical compositions were used to study the
Received 6 December 2006
hardenability of austenite with different alloy content and dispersion in dual-phase steel.
Received in revised form
The results showed that ferrite carbide aggregate did not form and considerable amount of
8 December 2007
the austenite present at the intercritical annealing temperatures transformed to martensite
Accepted 23 December 2007
even at the cooling rate of 0.01 ◦ C/s (furnace cooling) in the specimens. In dual-phase steels the absence of the carbide formation is unusual at this cooling rate and the critical cooling rates for the martensite formation is so low, compared to the previous studies. It is concluded
Keywords:
that without a high quenching power, these steel compositions are suitable for industrial
Dual-phase steels
production of dual-phase steel through continuous annealing line of iron and steel plants.
Alloy content
© 2008 Elsevier B.V. All rights reserved.
Phase transformation Hardenability
1.
Introduction
In order to improve the fuel economy in automotive industry, reduction in strength/weight ratio is required. For structural applications of automotive industry, a better combination of strength and ductility is the primary requirement in fabrication. These requirements were met by the advent of the dual-phase steels whose microstructure consist mainly of ferrite and martensite, exhibits low yield strength, high tensile strength, continuous yielding, and good uniform elongation in favour of forming applications. Commercially dual-phase sheet steels produced by intercritical heat treatment with either water quenching after continuous annealing (Matsudo et al., 1985; Pradhan, 1997) or box annealing (Pradhan, 1997; Osawa et al., 1978). Since 1975 most present production has concentrated on using continu-
∗
Corresponding author. Tel.: +90 312 2130494; fax: +90 312 2120059. ˘ E-mail address: [email protected] (M. Erdogan). 0924-0136/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2007.12.094
ous annealing because of their favour properties for satisfying sheets properties such as good surface, and optimum homogeneous properties all over of the sheets. The production of cold rolled dual-phase steels through continuous annealing process involves hot rolling, cold rolling, and continuous annealing in the (␣ + ␥)-phase field for the nucleation and the growth of austenite into ferrite, quenching for transformation of the austenite into martensite, overaging, and air cooling (Rocha et al., 2005). The degree of quenching power to obtain desired quantity of martensite depends on hardenability of austenite formed during intercritical annealing. From the economical point of view, production of cold rolled dual-phase steel through continuous annealing line, low quenching power is desirable, provided a useful quantity of martensite is obtained and pearlite and bainite are absent.
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In dual-phase heat treatment, the annealing temperatures in the (␣ + ␥)-phase field controls the volume fraction of austenite, thus affecting the hardenability of individual austenite pools. At the critical cooling rates, the austenite present fully transforms to martensite resulting in simple mix of ferrite and martensite. On slow cooling rates, the austenite pool first decreases in size by epitaxial ferrite growth on retained ferrite and at still slower rates, the remaining austenite enriched in carbon transforms to martensite, bainite and/or pearlite depending upon the alloy content of the austenite and the cooling rate. The data of constituent phases of microstructure obtained by the heat treatment in practice is routinely used in alloy design. The effect of varying heat treatment temperature, time, alloy content, and cooling rate on the decomposition product has been intensively studied by several investigators (Lawson et al., 1981; Eldis, 1979; Matlock et al., 1979; Lawson et al., 1980; Jeong and Kim, 1985; Marder, 1977; Geib et al., 1980; Mould and Skena, 1977; Rigsbee and Vander Arend, 1977; Huppi et al., 1980; Erdogan and Priestner, 1999; Erdogan and Priestner, 2002; Erdogan, 2002; Sarwar and Priestner, 1999; Priestner and Ajmal, 1987; Erdogan, 2003). The austenite transformation following intercritical annealing is different from that after normal austenitisation in two respects. Firstly in dual-phase heat treatment austenite volume fraction and its carbon content are determined by the intercritical annealing temperature. Under paraequilibrium condition (short intercritical annealing time) carbon, but not substitutional alloying elements segregates and determines phase proportions and compositions. Secondly, a nucleation step is not required for the new ferrite (it is also called epitaxial ferrite) to form during the cooling, because the old ferrite (proeutectoid ferrite) present during annealing can grow epitaxially into the austenite. Lawson et al. (1981) created microstructure maps which illustrated the transformed products quantitatively as a function of the cooling rate for a particular heat treatment temperature. They demonstrated that the total volume fraction of the transformed phases was approximately constant for all cooling rates. In addition to composition, intercritical annealing temperature and cooling rate affect hardenability of dual-phase steel. Priestner and Ajmal (1987) proposed that microstructure maps may be interpreted to obtain austenite → martensite hardenability data specific to different austenite contents. They plotted the percentage of the austenite which transformed to martensite as a function of cooling rate. Recently, Erdogan (2003) reported that ferrite carbide aggregate formation suppressed to cooling rates below 1 ◦ C/s and a significant amount of austenite transformed to martensite even at the cooling rate of 0.1 ◦ C/s in a steel containing
0.065 wt.% C, 1.58 wt.% Mn and 0.5 wt.% Ni. The critical cooling rates for carbide formation was so low, compared with previous studies. This result was attributed to the combination of low carbon content and the presence of Ni. Using the same steel it was also concluded that the fine dualphase structure produced more martensite than the coarse microstructure after annealing at low temperature and particular at the slower cooling rates. The reason for this was considered to the result from higher carbon enrichment effect in fine structure than the coarse ones during cooling. Based on Erdogan’s study it was considered that increment in Ni content with respect to Ni content in the previous study (Erdogan, 2003) may result the suppressing of ferrite carbide aggregate formation to lower cooling rates and a significant amount of austenite transforms to martensite at the lower cooling rates. This is important since low quenching power desirable for dual-phase production as mentioned above. In the current dual-phase literature, the effect of Ni content variation on austenite hardenability has not been reported yet. Therefore current investigation was designed to evaluate the effect of variation in particularly Ni and alloy content and austenite dispersion on austenite hardenability in dual-phase steels. For this purpose, two hot + cold rolled steels containing different chemical compositions and austenite dispersions (fine and course) at the different intercritical annealing temperatures were used to investigate austenite hardenability of these steels and a production process was followed to simulate the production process of cold rolled dual phase through industrial continuous annealing up to overaging stage.
2.
Experimental procedure
The steels were produced in a medium frequency vacuum induction furnace. Chemical compositions of used steels (steels 1 and 2) are given in Table 1. In fact, the study was designed to obtain two levels of nickel and keeping the other composition constant in both steels. However, following the two different steel production, some compositional differences in other composition were observed. As cast samples was supplied in the form of 300 mm × 300 mm × 26 mm plate. 26 × 26 × 300 square bars were cut from the plate and hot rolled to 4mm and finally cold rolled to 2 mm. After the cold rolling, 12 mm × 12 mm × 2 mm specimens from the cold rolled plates used for the simulation of the heating, soaking and cooling stages of industrial continuous annealing (Fig. 1). The preliminary investigation was to determine the dependence of martensite volume fraction (MVF) on intercritical annealing temperature (ICAT). For this purpose 12 mm × 12 mm × 2 mm specimens from the cold rolled plates
Table 1 – Chemical compositions of materials used in experiments Steels
1 2
Chemical composition (wt.%) C
Mn
Si
Ni
P
S
V
Ti
Nb
Fe
0.064 0.062
1.72 1.89
0.46 0.37
0.67 0.85
0.0113 0.0130
0.0105 0.0110
0.07 0.06
0.015 0.013
0.005 0.003
Rest
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Fig. 1 – Schematic diagram of the continuous annealing process.
respectively, for a detailed study of the effect of the cooling rate on the development of the dual-phase microstructure (Fig. 2). These temperatures were chosen in order to specify a series of heat treatments that would vary the new ferrite content at two levels of constant MVFs of ∼15 and ∼25% and two levels of microstructural refinement. Fine and coarse dispersions of martensite were obtained for steels 1 and 2 from two different starting conditions. Dualphase microstructures derived from as hot and then cold rolled starting microstructures were labeled “series A”. These materials had ferrite + pearlite structure (Fig. 3(a) and (b)). The other starting microstructure for the steels 1 and 2 were obtained by re-austenitising the series A specimens at 900 ◦ C for 20 min and the water quenching which produced a microstructure that was nearly wholly martensitic (Fig. 4(a) and (b)). Dual-phase microstructures derived from this initial microstructure were labeled “series B”. The microstructures of the specimens A and B were the starting point for subsequent intercritical annealing heat treatment.
2.1.
Fig. 2 – The dependence of austenite (martensite at room temperature) volume fraction on the intercritical annealing temperature.
were annealed for 20 min at a series of temperatures from approximately 707–875 ◦ C and from 703 to 870 ◦ C for Ae1 and Ae3 temperatures of steels 1 and 2, respectively (Fig. 2) and then quenched into water held at room temperature. The Ac1 and Ac3 temperature limits were calculated theoretically from the chemical composition of the material (Rocha et al., 2005; Andrew, 1965; Hillis et al., 1998; Basso et al., 2007; Kaplan et al., 2007). The MVF was determined by the point counting on metallographic sections etched in 2% nital. As a result of this preliminary study, three ICATs of 717, 737, 795 ◦ C and 735, 755, 815 ◦ C were selected for steels 1 and 2,
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Determination of microstructure maps
In order to control the heat treatment of the specimens in a manner that would vary proportions of martensite and epitaxial ferrite, it was necessary to construct the microstructure maps for the two starting microstructures. Based on the results of the preliminary investigation, 12 mm × 12 mm × 2 mm specimens that had been previously heat treated to produce the starting microstructures, A and B, were intercritically annealed at 717, 737, 795 ◦ C and 735, 755, 815 ◦ C, respectively, for 20 min. They were then cooled to room temperature in range of cooling media which resulted in average cooling rates over the first half of the cooling range of about 1500–0.01 ◦ C/s. The proportions of the constituents then present were determined by point counting on etched metallographic sections are listed in the Tables 3 and 4. Between 1000 and 2000 points were counted, to keep the standard error below 6% of the smallest phase volume fraction present in a sample. Typical cooling rates are listed in Table 2. Actual cooling rates were used in plotting data in the form of microstructure maps (see Section 3.2). Throughout these heat treatments, the temperature of each specimen was monitored by a thermocouple spot-welded to the center of one of its faces.
Fig. 3 – The starting microstructure of series A specimens for steel 1 (a) and steel 2 (b), etchant: 2% nital.
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Fig. 4 – The starting microstructure of series B specimens for steel 1 (a) and steel 2 (b), etchant: 2% nital.
Table 2 – Cooling rates and phase proportions in steel 1 specimens Specimen code 1A735(15) 1A735(25) 1B735(15) 1B735(25) 1A755(15) 1A755(25) 1B755(15) 1B755(25) 1A815(15) 1A815(25) 1B815(15) 1B815(25)
Cooling rate (◦ C/s)
Martensite content (vol.%)
0.2 1400 0.3 1400 0.34 5 0.041 4.5 0.56 80 0.45 100
12 25 17 25 17 24 16 26.5 14.6 25 14 27
The amount of retained austenite in the specimens was measured by X-ray diffractometry following the standard procedure in which the ratio of the area under austenite and ferrite peaks was converted into volume fractions (Cullity, 1978). The analysis was carried out on the surface of the specimens after grinding them and polishing with 1 m diamond paste. X-ray analysis used a Cu K␣ source to identify the austenite-phase volume fraction. This analysis failed to detect retained austenite in all but a small number of the specimens sampled, and the retained austenite was present only at the limit of the methods sensitivity <3% by volume, in the exceptions. The use of XRD for the quantitative evaluation of the austenite is not so simple and is a still a matter of scientific discussions (Zhao et al., 2001). The mean linear intercept grain size of the ferrite matrix was estimated by the superimposing circles on the micrographs at a magnification that allowed at least 500 intercepts to be counted and standard deviation ( d ) calculated by d = √ grain size × 0.7/ 500 and adopted to ASTM grain size number. The average grain size and its standard deviation in the A series of steels 1 and 2 specimens were 7 ± 0.21 m (ASTM 5) and 10 ± 0.32 (ASTM 6), respectively, and that in the B series for steels 1 and 2 were 4 ± 0.13 m (ASTM 4) and 6 ± 0.18 (ASTM 5), respectively. Results showed that B series have finer ferrite grains than A series. Steel 2 also showed finer grains compared to the steel 1. The specimens were coded according to steel number, starting microstructure, ICAT and nominal MVF. For exam-
New ferrite content (vol.%) 14 – 8 – 18 12 20 9.5 40.4 30 41 28
ple, in specimen code 1B735(15), 1 stands for steel number, B for the starting microstructure, 735 for ICAT and (15) for MVF.
3.
Results and discussion
3.1.
Microstructures
On heating starting microstructure A to the ICAT, austenite (martensite at room temperature) nucleated at, and grew to consume, the pearlite and then grew into the ferrite. The quenching from different ICAT produced martensitic structure with different MVF restricted to ferrite/ferrite boundaries and an isolated or continuous network of martensite structure along ferrite/ferrite boundaries formed depending on MVFs (Fig. 5(a) and (b)). During intercritical annealing of starting microstructure B, the preferential area for austenite nucleation was the lath interfaces, lath colony boundaries and prior austenite grain boundaries. The austenite volumes obtained on intercritical annealing and then water quenching was then smaller and more finely dispersed (Fig. 6(a) and (b)). This result is in good agreement with literature (Erdogan and Priestner, 1999, 2002; Erdogan, 2002; Hughes, 1986; Jamiru, 1990; Thomas and Koo, 1981; Tomota, 1987). The size and distribution of the martensite obtained on quenching then reflected the scale of the initial microstructure.
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Fig. 5 – Micrograph of 1A735(15) (a) and 2A755(25) (b) specimens. 1A735(15) specimen, intercritically annealed at 735 ◦ C and cooled at 0.3 ◦ C/s to give 17% martensite and 8% new ferrite and 2A755(25) specimen, intercritically annealed at 737 ◦ C and cooled at 5 ◦ C/s to give 25% martensite and 10% new ferrite etched in 2% nital.
Fig. 6 – Micrograph of 2B717(15) (a) and 2B717(25) (b) specimens. 2B717(15) specimen, intercritically annealed at 717 ◦ C and cooled at 0.3 ◦ C/s to give 14% martensite and 12% new ferrite and 2B717(25) specimen, intercritically annealed at 737 ◦ C and cooled at 1400 ◦ C/s to give 26% martensite and 0% new ferrite etched in 2% nital.
Fig. 7 – Quantitative microstructure maps of series A specimens of steel 1 (a) and steel 2 (b) annealed at 735 ◦ C showing effect of cooling rates on the microstructure (old ferrite: ferrite present during intercritically annealing).
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3.2.
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Quantitative microstructure maps
Figs. 7 and 8 illustrate some examples of the effect of cooling rate on the proportion of phases after heat treatment in the form of the microstructure maps for series A and B starting microstructures. Since the martensite and new ferrite transforms from austenite present at ICAT during cooling, the sum of their volume fraction should be equal to the volume fraction of austenite present at the intercritical temperature (although some austenite may be retained in some dual-phase steels). For the present steel, the cooling rates in excess of 1000 ◦ C /s were required to convert all of the austenite into martensite. With decreasing the cooling rates, martensite was increasingly replaced by the new ferrite. At the slower cooling rates, the formation of the new ferrite was followed by the transformation of remaining austenite to martensite and carbide was absent. Some retained austenite was probably present at some cooling rates, as will be discussed later. At any cooling rates, the absence of the ferrite–carbide aggregate formation is an important feature of the transformation characteristics in these steels since the presence of carbide in dual-phase steels has a detrimental effect on ductility. The other important feature of those microstructure maps (Figs. 7 and 8) at the cooling rate of ∼50 ◦ C/s corresponding to the required cooling rate for industrial application of rapid cooling stage IV (Fig. 1), carbide is absent and MVF is over 20%. The absence of carbide formation is unusual at this cooling rate and the critical cooling rates for martensite formation compared with the previous studies is so low. Generally dualphase steels having MVF between 15 and 25% exhibit optimum mechanical properties (Rocha et al., 2005). Some examples of microstructures obtained after cooling from ICAT ranges at different cooling rates are shown in Fig. 9(a) and (b). Cooling the specimen 2B717(15) from 717 ◦ C (Fig. 9(a)) converted most of the austenite to martensite, the
martensite volume fraction being close to ∼15% and the new ferrite content being ∼12%. Cooling the specimen 2B795(15) from the temperature of 795 ◦ C (Fig. 9(b)), resulted in more new ferrite and 17% martensite. In Fig. 9(a) and (b), the chromate etch invented by Lawson et al. (1980) was used to show the new ferrite as the bright, white constituent in the micrographs, 12% in Fig. 9(a) and 39% in Fig. 9(b). The critical cooling rate for martensite formation is quite low when compared with the previous studies; (Lawson et al., 1980, 1981; Eldis, 1979; Matlock et al., 1979; Jeong and Kim, 1985; Marder, 1977; Geib et al., 1980; Mould and Skena, 1977; Rigsbee and Vander Arend, 1977; Huppi et al., 1980). The reason for this may be due to the combination of low carbon and the high Ni contents. As is mentioned before, in the current dual-phase literature, effect of Ni content variation on austenite hardenability has not been reported yet. Mn is well known to have an effect on austenite hardenability and is without any exception present in dual-phase steels with changing amounts in different alloys. However, in the previous studies (Lawson et al., 1980, 1981; Eldis, 1979; Matlock et al., 1979; Jeong and Kim, 1985; Marder, 1977; Geib et al., 1980; Mould and Skena, 1977; Rigsbee and Vander Arend, 1977; Huppi et al., 1980; Erdogan and Priestner, 1999, 2002; Erdogan, 2002; Sarwar and Priestner, 1999), still high amount of ferrite carbide aggregate have been observed to be present even in dual-phase steels which have high Mn contents at slower cooling rates. Therefore, the effect of Ni on austenite hardenability is focused on and underlined here. Figs. 7 and 8 also show that a significant amount of austenite is transformed to martensite even at the cooling rate of 0.01 ◦ C/s (furnace cooling). At the same MVF, much less new ferrite was present after cooling from low intercritical annealing temperatures than after cooling from the highest intercritical annealing temperature (Tables 3 and 4). This is due to the higher carbon concentration in austenite at lower ICAT and consequently higher hardenability 19].
Fig. 8 – Quantitative microstructure maps of series B specimens of steel 1 (a) and steel 2 (b) annealed at 735 ◦ C showing effect of cooling rates on the microstructure (old ferrite: ferrite present during intercritically annealing).
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Fig. 9 – Micrograph of 2B717(15) (a) and 2B795(15) (b) specimens. 2B717(15) specimen, intercritically annealed at 717 ◦ C and cooled at 0.3 ◦ C/s to give 14% martensite and 12% new ferrite and 2B795(15) specimen, intercritically annealed at 795 ◦ C and cooled at 0.54 ◦ C /s to give 17% martensite and 39% new ferrite. New ferrite revealed as bright white areas after etching in hot chromate reagent.
3.3.
Effect of austenite dispersion on its hardenability
The microstructure maps shown in Figs. 7 and 8 were interpreted in the form of austenite → martensite hardenability diagrams (Priestner and Ajmal, 1987), as in Fig. 10(a) and (b). In dual-phase steels, the proportions of the several constituents present after dual-phase heat treatment depend on the cooling rate, the composition of the austenite and the degree of segregation of the substitutional element prior to the cooling from the ICAT. In addition, Priestner (1987) proposed that the dispersion of the austenite in ferrite at the ICAT should have a significant effect on the hardenability. The dispersion of the austenite particles in the ferrite was characterized by the ratio of interfacial area to austenite volume fraction. Priestner showed by geometric analysis that as this ratio increased the yield of martensite at any cooling rate should decrease, because for a given rate of epitaxial ferrite growth, a large interfacial area would result in a larger volume of new ferrite when the temperature fell to the Ms of
the austenite. For a given volume fraction of austenite, a high interfacial area equates to fine austenite particles. Hughes (1986) demonstrated that the starting microstructure defined the morphology of the austenite at the intercritical annealing temperature. His study was on 0.12% C, 1.5% Mn, 0.34% Si steel, which was homogenized at 1150 ◦ C for 20 h and then heat treated to provide starting microstructures of different scales. On intercritical annealing, the different starting microstructures produced a range of austenite dispersions. The finer starting microstructure gave a finer, more dispersed austenite, which needed faster cooling to achieve a constant volume fraction of martensite, nearly the same with Priestner’s proposal. Using the same steel, Jamiru (1990) also investigated the effect of austenite dispersion on its hardenability. From the different starting microstructures a range of austenite dispersions were yielded by intercritical annealing. He also showed that there was a strong reduction in martensite yield as the dispersion parameter increased for all cooling rates. However,
Table 3 – Cooling rates and phase proportions in steel 2 specimens Specimen code 2A717(15) 2A717(25) 2B717(15) 2B717(25) 2A737(15) 2A737(25) 2B737(15) 2B737(25) 2A795(15) 2A795(25) 2B795(15) 2B795(25)
Cooling rate (◦ C/s)
Martensite content (vol.%)
0.3 1400 0.042 1400 0.04 5 0.04 5 0.54 50 0.52 50
14 26 14 26 15.5 25 17 26 17 27 15 26
New ferrite content (vol.%) 12 – 12 – 19.5 10 18 9 39 29 41 30
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Fig. 10 – Martensite hardenability of various austenites of steels.
he pointed out that at the slower cooling rates the relation between the dispersion and the martensite content was weaker than the geometric model predicted. The reason for the disagreement between the model and the experimental data was explained as follows; in the model, it was assumed that the cooling rate was constant and that an average ferrite growth velocity applied throughout cooling. In reality this is not the case. Ferrite growth rate is a complex function of the degree of undercooling, carbon diffusion rate and carbon concentration in the austenite at the ␣/␥ interface. The model ignores the fact that small particles of austenite must become enriched in carbon more rapidly than large particles, with respect to the amount of ferrite growth. This suggests that for a given cooling rate, the average ferrite growth rate will be considerably slower for fine dispersions than the coarse ones. Therefore, less ferrite and more martensite will form relative to the geometric theory, and this difference will increase with increasing austenite dispersion. In the present experiments the finer starting microstructure, series B, produced more martensite than the coarser A, after the intercritical heat treatments at the lower temperatures, and particularly at the slower cooling rates, i.e. at the slowest cooling rate of 0.01 ◦ C/s, both MVFs of finer B and coarse A series for steels 1 and 2 are ∼11% (Fig. 8(a)) and ∼7% (Fig. 7(a)) respectively. Under the same cooling condition, finer series B of steel 2 (Fig. 8(b)) were also produced more MVF than coarse series A of steel 2 (Fig. 8(b). After the intercritical heat treatment at the highest temperature, the hardenability of austenite from both series was nearly identical. To explain these results, it is necessary to propose that the carbon-enrichment effect is capable of balancing or overwhelming Priestner’s geometric effect. In the present work, the concentration of the austenite stabilizing elements (1.72–1.89% Mn and 0.67–0.85% Ni) is higher than in Jamiru’s and Erdogan’s (Erdogan, 2003) steel, and would depress the Ms temperature. Also, the carbon content of the present steel is only half that of Jamiru’s steel and almost same as Erdogan’s steel. In fully austenitised steel, the critical cooling rate for 100% martensite is increased by reducing
the carbon content and decreased by increasing the substitutional alloy content. In the present steels, this critical cooling rate after the intercritical heat treatment is approximately 1000 ◦ C/s, which is similar to that for higher carbon, lower alloyed steels investigated previously. However, a substantially larger fraction of martensite is retained to much lower cooling rates in the present steels. This suggests that the alloying elements have a large effect in reducing the ferrite growth rate, despite the lower carbon concentration. However, it will be noted that the ferrite growth must be very sensitive to carbon concentration in the austenite. Therefore, much more new ferrite has been formed in the austenite obtained at high ICAT than the one at low ICAT ones. It is the sensitivity of ferrite growth rate to the increasing carbon concentration in the austenite as the austenite shrinks during the cooling. That negates Priestner’s geometric effect when the materials have been annealed at the high ICAT, but after annealing at the low ICAT, Priestner’s theory has been seen to be approved. When both steels have been cooled from approximately the same temperature of 735 or 737 ◦ C, the required cooling rates to obtain 25% martensite are 1400 ◦ C/s for the steel 1 and only 5 ◦ C/s for the steel 2 (Figs. 7 and 8). The influence of Mn and Ni on the hardenability by depressing the Ms temperature is well known. The amount of Mn and Ni in the steel 2 is higher than the steel 1 and both steels have the same carbon content. In addition to martensite and ferrite, retained austenite may be present in dual-phase steel in amounts varying from 2 to 9% (Rigsbee and Vander Arend, 1977; Yi and Kim, 1983; Jeong and Kim, 1987; Bangaru and Sachdev, 1982; Furukawa et al., 1979; Saleh and Priestner, 2001). The amount of retained austenite is sensitive to the cooling rate, intercritical annealing temperature and carbon content of the steel. Rigsbee and Vander Arend (Rigsbee and Vander Arend, 1977) and Bangaru and Sachdev (Bangaru and Sachdev, 1982) reported that the retained austenite content increased slightly with the increasing ICAT and the carbon content of the steel. It has been observed (Priestner, 1987; Yi and Kim, 1983; Jeong and Kim, 1987) that the amount of the retained austenite also increased as the cooling rate decreased. In the present work the retained
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austenite was found to be barely detectable in some samples cooled slowly.
4.
Conclusions
The following conclusions could be drawn from the above discussions. 1. The critical cooling rate for the martensite formation is considerable low compared with the previous studies. This is probably due to the combination of low carbon and high Ni contents. 2. In the present steels ferrite-carbide aggregates did not form at the regardless of any cooling rates. This result is unusual compared with the previous studies. 3. The fine dual-phase structure produced more martensite than the coarse microstructure after annealing at the low temperature and particular at the slower cooling rates. The reason for this was considered to the result from higher carbon enrichment effect in the fine structure than the coarse ones during the cooling. 4. With considerable low quenching power, the compositions of the present steels are suitable for industrial production of dual-phase steel through continuous annealing line of iron and steel plants, providing a useful quantity of martensite and pearlite, and ferrite-carbide aggregates are absent.
Acknowledgements The authors wish to acknowledge the financial supports of Eregli Iron and Steel Plant, Turkey for this study. The author is also indebted to The Scientific and Technical Research Council of Turkey, High Technology Institute Gebze, Turkey for providing their induction furnace castings and cold rolling facilities.
references
Andrew, K.W., 1965. Empirical formulae for the calculation of some transformation temperatures. J. Iron Steel Inst. 203, 721–727. Bangaru, N.R.V., Sachdev, A.K., 1982. Influence of cooling rate on the microstructure and retained austenite in an intercritical annealed vanadium containing HSLA steel. Met. Trans. A 13A, 1899–1906. Basso, A.D., Mart´ınez, R.A., Sikora, J.A., 2007. Influence of austenitising and austempering temperatures on microstructure and properties of dual phase ADI. Mater. Sci. Technol. 23 (11), 1321–1326. Cullity, B.D., 1978. Elements of X-Ray Diffraction, 2nd edn. Addison-Wesley Series in Metallurgy and Materials, USA. Eldis, G.T., 1979. Influence of microstructure and testing procedure on the measured mechanical properties of heat-treated dual-phase steels. In: Kott, R.A., Morris, J.W. (Eds.), Structure and Properties of Dual-Phase Steels. AIME, New York, pp. 202–220. Erdogan, M., 2002. The effect of new ferrite content on the tensile fracture behaviour of dual phase steels. J. Mater. Sci. 37, 3623–3630. Erdogan, M., 2003. The effect of austenite dispersion on phase transformation in dual phase steels. Scr. Mater. 48, 501–506.
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Erdogan, M., Priestner, R., 1999. Effect of epitaxial ferrite on yielding and plastic flow in dual phase steel in tension and compression. Mater. Sci. Technol. 15, 1273–1284. Erdogan, M., Priestner, R., 2002. Effect of martensite content, its dispersion, and epitaxial ferrite content on Bauschinger behaviour of dual phase steel. Mater. Sci. Technol. 18, 369–376. Furukawa, T., Morikawa, H., Takechi, H., Koyama, K., 1979. Process factors for highly ductile dual-phase sheet steels. In: Kott, R.A., Morris, J.W. (Eds.), Structure and Properties of Dual-Phase Steels. AIME, New York, pp. 281–303. Geib, M.D., Matlock, D.K., Krauss, G., 1980. The effect of intercritical temperature on the structure of niobium microalloyed dual-phase steel. Met. Trans. A 11A, 1683–1689. Hillis, D.J., Llewellyn, D.T., Evans, P.J., 1998. Rapid annealing of dual phase steels. Ironmaking Steelmaking 25 (1), 47–54. G. Hughes, The influence of prior microstructure on the constitution of dual-phase steels, M.Sc. Thesis, University of Manchester, Manchester, UK, May 1986. Huppi, G.S., Matlock, D.K., Krauss, G., 1980. An evaluation of the importance of epitaxial ferrite in dual-phase steel microstructures. Scr. Met. 25, 1239–1243. T. Jamiru, Effect of austenite dispersion in dual phase steel on its hardenability, M.Sc. Thesis, University of Manchester, 1990. Jeong, W.C., Kim, C.H., 1985. Ferrite growth on cooling after intercritical annealing in a dual phase steel. Scr. Met. 118, 37–42. Jeong, W.C., Kim, C.H., 1987. On the isolated retained austenite in dual phase steel. Met. Trans. A 18A, 933–934. ¨ C.H., Erdogan, M., 2007. Characterization of Kaplan, M., Gur, dual-phase steels using magnetic Barkhausen noise technique. J. Nondestruct. Eval. 26, 79–87. Lawson, R.D., Matlock, D.K., Krauss, G., 1980. Etching technique for microalloyed dual-phase steels. Metallography 13, 71–87. Lawson, R.D., Matlock, D.K., Krauss, G., 1981. The effect of microstructure on the deformation behaviour and mechanical properties of a dual phase steel. In: Kott, R.A., Bramfitt, B.L. (Eds.), Fundamentals of Dual-Phase Steels. AIME, New York, pp. 347–381. Marder, A.R., 1977. Factors affecting the ductility of dual phase steels. In: Davenport, A.T. (Ed.), Formable HSLA and Dual Phase Steels. AIME, New York, pp. 87–98. Matlock, D.K., Krauss, G., Ramos, L.F., Huppi, G.S., 1979. A correlation of processing variables with deformation of dual-phase steels. In: Kott, R.A., Morris, J.W. (Eds.), Structure and Properties of Dual-Phase Steels,. AIME, New York, pp. 62–89. Matsudo, K., Osuma, K., Kurikava, K., 1985. Metallurgical aspects of the development continuous annealing technology at Nippin Kokan. Metall. Soc. AIME, 3–36. Mould, P.R., Skena, C.C., 1977. Structure and properties of cold rolled ferrite phase martensite (dual-phase) steels sheets. In: Davenport, A.T. (Ed.), Formable HSLA and Dual Phase Steels. AIME, New York, pp. 183–203. Osawa, K., Shimomura, T., Kinoshita, M., Matsudo, K., Iwase, K., 1978. Development of high strength cold rolled steel sheets for automotive use by continuous annealing. J. Met. 30, 6–9. Pradhan, R., 1997. Continuous annealing of the steel. ASM Handbook 4, 56–66 (5th printing). Priestner, R., 1987. Effect of austenite dispersion and plastic deformation on phase transformation in dual phase steel. In: Conf. Proc. Phase Transformation ’87. Institute of Metals, Cambridge, pp. 411–413. Priestner, R., Ajmal, M., 1987. Effect of carbon content and microalloying on martensitic hardenability of austenite of dual phase steel. Mater. Sci. Technol. 3, 360–364. Rigsbee, J.M., Vander Arend, P.J., 1977. Laboratory studies of microstructures and structure–property relationships in dual-phase steels. In: Davenport, A.T. (Ed.), Formable HSLA and Dual Phase Steels. AIME, New York, pp. 56–85.
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Rocha, R.O., Melo, T.M.F., Pereloma, E.V., Santos, D.B., 2005. Microstructural evolution at the initial stages of continuous annealing of cold rolled dual-phase steel. Mater. Sci. Eng. A 391, 296–304. Saleh, M.H., Priestner, R., 2001. Retained austenite in dual phase silicon steels and effect on mechanical properties. J. Mater. Process. Technol. 113, 587–593. Sarwar, M., Priestner, R., 1999. Hardenability of austenite in a dual-phase steel. J. Mater. Eng. Perform. 83, 380. Thomas, G., Koo, Y.J., 1981. Developments in strong ductile duplex ferritic–martensitic steels. In: Kott, R.A., Bramfitt, B.L.
(Eds.), Fundamentals of Dual-Phase Steels. AIME, New York, pp. 183–201. Tomota, Y., 1987. Effects of morphology and strength of martensite on cyclic deformation behaviour in dual-phase steels. Mater. Sci. Technol. 3, 415–421. Yi, J.J., Kim, I.S., 1983. Ferrite growth during slow cooling from the intercritical annealing of a dual phase steel. Scr. Met. 17, 299–302. ¨ Zhao, L., Van Dijk, N.H., Bruck, E., Sietsma, J., Van Der Zwaag, S., 2001. Magnetic and X-ray diffraction measurements for the determination of retained austenite in TRIP steels. Mater. Sci. Eng. A 313, 145–153.
journal homepage: www.elsevier.com/locate/jmatprotec
The hardenability of austenite with different alloy content and dispersion in dual-phase steels ˘ b,∗ B. Demir a , M. Erdogan a
Materials Division, Metallurgy Education Department, Faculty of Technical Education, Karaelmas University, 78050 Karabuk, Turkey b Materials Division, Metallurgy Education Department, Faculty of Technical Education, Gazi University, 06500 Besevler, Ankara, Turkey
a r t i c l e
i n f o
a b s t r a c t
Article history:
In this investigation, two steels with different chemical compositions were used to study the
Received 6 December 2006
hardenability of austenite with different alloy content and dispersion in dual-phase steel.
Received in revised form
The results showed that ferrite carbide aggregate did not form and considerable amount of
8 December 2007
the austenite present at the intercritical annealing temperatures transformed to martensite
Accepted 23 December 2007
even at the cooling rate of 0.01 ◦ C/s (furnace cooling) in the specimens. In dual-phase steels the absence of the carbide formation is unusual at this cooling rate and the critical cooling rates for the martensite formation is so low, compared to the previous studies. It is concluded
Keywords:
that without a high quenching power, these steel compositions are suitable for industrial
Dual-phase steels
production of dual-phase steel through continuous annealing line of iron and steel plants.
Alloy content
© 2008 Elsevier B.V. All rights reserved.
Phase transformation Hardenability
1.
Introduction
In order to improve the fuel economy in automotive industry, reduction in strength/weight ratio is required. For structural applications of automotive industry, a better combination of strength and ductility is the primary requirement in fabrication. These requirements were met by the advent of the dual-phase steels whose microstructure consist mainly of ferrite and martensite, exhibits low yield strength, high tensile strength, continuous yielding, and good uniform elongation in favour of forming applications. Commercially dual-phase sheet steels produced by intercritical heat treatment with either water quenching after continuous annealing (Matsudo et al., 1985; Pradhan, 1997) or box annealing (Pradhan, 1997; Osawa et al., 1978). Since 1975 most present production has concentrated on using continu-
∗
Corresponding author. Tel.: +90 312 2130494; fax: +90 312 2120059. ˘ E-mail address: [email protected] (M. Erdogan). 0924-0136/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2007.12.094
ous annealing because of their favour properties for satisfying sheets properties such as good surface, and optimum homogeneous properties all over of the sheets. The production of cold rolled dual-phase steels through continuous annealing process involves hot rolling, cold rolling, and continuous annealing in the (␣ + ␥)-phase field for the nucleation and the growth of austenite into ferrite, quenching for transformation of the austenite into martensite, overaging, and air cooling (Rocha et al., 2005). The degree of quenching power to obtain desired quantity of martensite depends on hardenability of austenite formed during intercritical annealing. From the economical point of view, production of cold rolled dual-phase steel through continuous annealing line, low quenching power is desirable, provided a useful quantity of martensite is obtained and pearlite and bainite are absent.
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In dual-phase heat treatment, the annealing temperatures in the (␣ + ␥)-phase field controls the volume fraction of austenite, thus affecting the hardenability of individual austenite pools. At the critical cooling rates, the austenite present fully transforms to martensite resulting in simple mix of ferrite and martensite. On slow cooling rates, the austenite pool first decreases in size by epitaxial ferrite growth on retained ferrite and at still slower rates, the remaining austenite enriched in carbon transforms to martensite, bainite and/or pearlite depending upon the alloy content of the austenite and the cooling rate. The data of constituent phases of microstructure obtained by the heat treatment in practice is routinely used in alloy design. The effect of varying heat treatment temperature, time, alloy content, and cooling rate on the decomposition product has been intensively studied by several investigators (Lawson et al., 1981; Eldis, 1979; Matlock et al., 1979; Lawson et al., 1980; Jeong and Kim, 1985; Marder, 1977; Geib et al., 1980; Mould and Skena, 1977; Rigsbee and Vander Arend, 1977; Huppi et al., 1980; Erdogan and Priestner, 1999; Erdogan and Priestner, 2002; Erdogan, 2002; Sarwar and Priestner, 1999; Priestner and Ajmal, 1987; Erdogan, 2003). The austenite transformation following intercritical annealing is different from that after normal austenitisation in two respects. Firstly in dual-phase heat treatment austenite volume fraction and its carbon content are determined by the intercritical annealing temperature. Under paraequilibrium condition (short intercritical annealing time) carbon, but not substitutional alloying elements segregates and determines phase proportions and compositions. Secondly, a nucleation step is not required for the new ferrite (it is also called epitaxial ferrite) to form during the cooling, because the old ferrite (proeutectoid ferrite) present during annealing can grow epitaxially into the austenite. Lawson et al. (1981) created microstructure maps which illustrated the transformed products quantitatively as a function of the cooling rate for a particular heat treatment temperature. They demonstrated that the total volume fraction of the transformed phases was approximately constant for all cooling rates. In addition to composition, intercritical annealing temperature and cooling rate affect hardenability of dual-phase steel. Priestner and Ajmal (1987) proposed that microstructure maps may be interpreted to obtain austenite → martensite hardenability data specific to different austenite contents. They plotted the percentage of the austenite which transformed to martensite as a function of cooling rate. Recently, Erdogan (2003) reported that ferrite carbide aggregate formation suppressed to cooling rates below 1 ◦ C/s and a significant amount of austenite transformed to martensite even at the cooling rate of 0.1 ◦ C/s in a steel containing
0.065 wt.% C, 1.58 wt.% Mn and 0.5 wt.% Ni. The critical cooling rates for carbide formation was so low, compared with previous studies. This result was attributed to the combination of low carbon content and the presence of Ni. Using the same steel it was also concluded that the fine dualphase structure produced more martensite than the coarse microstructure after annealing at low temperature and particular at the slower cooling rates. The reason for this was considered to the result from higher carbon enrichment effect in fine structure than the coarse ones during cooling. Based on Erdogan’s study it was considered that increment in Ni content with respect to Ni content in the previous study (Erdogan, 2003) may result the suppressing of ferrite carbide aggregate formation to lower cooling rates and a significant amount of austenite transforms to martensite at the lower cooling rates. This is important since low quenching power desirable for dual-phase production as mentioned above. In the current dual-phase literature, the effect of Ni content variation on austenite hardenability has not been reported yet. Therefore current investigation was designed to evaluate the effect of variation in particularly Ni and alloy content and austenite dispersion on austenite hardenability in dual-phase steels. For this purpose, two hot + cold rolled steels containing different chemical compositions and austenite dispersions (fine and course) at the different intercritical annealing temperatures were used to investigate austenite hardenability of these steels and a production process was followed to simulate the production process of cold rolled dual phase through industrial continuous annealing up to overaging stage.
2.
Experimental procedure
The steels were produced in a medium frequency vacuum induction furnace. Chemical compositions of used steels (steels 1 and 2) are given in Table 1. In fact, the study was designed to obtain two levels of nickel and keeping the other composition constant in both steels. However, following the two different steel production, some compositional differences in other composition were observed. As cast samples was supplied in the form of 300 mm × 300 mm × 26 mm plate. 26 × 26 × 300 square bars were cut from the plate and hot rolled to 4mm and finally cold rolled to 2 mm. After the cold rolling, 12 mm × 12 mm × 2 mm specimens from the cold rolled plates used for the simulation of the heating, soaking and cooling stages of industrial continuous annealing (Fig. 1). The preliminary investigation was to determine the dependence of martensite volume fraction (MVF) on intercritical annealing temperature (ICAT). For this purpose 12 mm × 12 mm × 2 mm specimens from the cold rolled plates
Table 1 – Chemical compositions of materials used in experiments Steels
1 2
Chemical composition (wt.%) C
Mn
Si
Ni
P
S
V
Ti
Nb
Fe
0.064 0.062
1.72 1.89
0.46 0.37
0.67 0.85
0.0113 0.0130
0.0105 0.0110
0.07 0.06
0.015 0.013
0.005 0.003
Rest
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Fig. 1 – Schematic diagram of the continuous annealing process.
respectively, for a detailed study of the effect of the cooling rate on the development of the dual-phase microstructure (Fig. 2). These temperatures were chosen in order to specify a series of heat treatments that would vary the new ferrite content at two levels of constant MVFs of ∼15 and ∼25% and two levels of microstructural refinement. Fine and coarse dispersions of martensite were obtained for steels 1 and 2 from two different starting conditions. Dualphase microstructures derived from as hot and then cold rolled starting microstructures were labeled “series A”. These materials had ferrite + pearlite structure (Fig. 3(a) and (b)). The other starting microstructure for the steels 1 and 2 were obtained by re-austenitising the series A specimens at 900 ◦ C for 20 min and the water quenching which produced a microstructure that was nearly wholly martensitic (Fig. 4(a) and (b)). Dual-phase microstructures derived from this initial microstructure were labeled “series B”. The microstructures of the specimens A and B were the starting point for subsequent intercritical annealing heat treatment.
2.1.
Fig. 2 – The dependence of austenite (martensite at room temperature) volume fraction on the intercritical annealing temperature.
were annealed for 20 min at a series of temperatures from approximately 707–875 ◦ C and from 703 to 870 ◦ C for Ae1 and Ae3 temperatures of steels 1 and 2, respectively (Fig. 2) and then quenched into water held at room temperature. The Ac1 and Ac3 temperature limits were calculated theoretically from the chemical composition of the material (Rocha et al., 2005; Andrew, 1965; Hillis et al., 1998; Basso et al., 2007; Kaplan et al., 2007). The MVF was determined by the point counting on metallographic sections etched in 2% nital. As a result of this preliminary study, three ICATs of 717, 737, 795 ◦ C and 735, 755, 815 ◦ C were selected for steels 1 and 2,
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Determination of microstructure maps
In order to control the heat treatment of the specimens in a manner that would vary proportions of martensite and epitaxial ferrite, it was necessary to construct the microstructure maps for the two starting microstructures. Based on the results of the preliminary investigation, 12 mm × 12 mm × 2 mm specimens that had been previously heat treated to produce the starting microstructures, A and B, were intercritically annealed at 717, 737, 795 ◦ C and 735, 755, 815 ◦ C, respectively, for 20 min. They were then cooled to room temperature in range of cooling media which resulted in average cooling rates over the first half of the cooling range of about 1500–0.01 ◦ C/s. The proportions of the constituents then present were determined by point counting on etched metallographic sections are listed in the Tables 3 and 4. Between 1000 and 2000 points were counted, to keep the standard error below 6% of the smallest phase volume fraction present in a sample. Typical cooling rates are listed in Table 2. Actual cooling rates were used in plotting data in the form of microstructure maps (see Section 3.2). Throughout these heat treatments, the temperature of each specimen was monitored by a thermocouple spot-welded to the center of one of its faces.
Fig. 3 – The starting microstructure of series A specimens for steel 1 (a) and steel 2 (b), etchant: 2% nital.
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Fig. 4 – The starting microstructure of series B specimens for steel 1 (a) and steel 2 (b), etchant: 2% nital.
Table 2 – Cooling rates and phase proportions in steel 1 specimens Specimen code 1A735(15) 1A735(25) 1B735(15) 1B735(25) 1A755(15) 1A755(25) 1B755(15) 1B755(25) 1A815(15) 1A815(25) 1B815(15) 1B815(25)
Cooling rate (◦ C/s)
Martensite content (vol.%)
0.2 1400 0.3 1400 0.34 5 0.041 4.5 0.56 80 0.45 100
12 25 17 25 17 24 16 26.5 14.6 25 14 27
The amount of retained austenite in the specimens was measured by X-ray diffractometry following the standard procedure in which the ratio of the area under austenite and ferrite peaks was converted into volume fractions (Cullity, 1978). The analysis was carried out on the surface of the specimens after grinding them and polishing with 1 m diamond paste. X-ray analysis used a Cu K␣ source to identify the austenite-phase volume fraction. This analysis failed to detect retained austenite in all but a small number of the specimens sampled, and the retained austenite was present only at the limit of the methods sensitivity <3% by volume, in the exceptions. The use of XRD for the quantitative evaluation of the austenite is not so simple and is a still a matter of scientific discussions (Zhao et al., 2001). The mean linear intercept grain size of the ferrite matrix was estimated by the superimposing circles on the micrographs at a magnification that allowed at least 500 intercepts to be counted and standard deviation ( d ) calculated by d = √ grain size × 0.7/ 500 and adopted to ASTM grain size number. The average grain size and its standard deviation in the A series of steels 1 and 2 specimens were 7 ± 0.21 m (ASTM 5) and 10 ± 0.32 (ASTM 6), respectively, and that in the B series for steels 1 and 2 were 4 ± 0.13 m (ASTM 4) and 6 ± 0.18 (ASTM 5), respectively. Results showed that B series have finer ferrite grains than A series. Steel 2 also showed finer grains compared to the steel 1. The specimens were coded according to steel number, starting microstructure, ICAT and nominal MVF. For exam-
New ferrite content (vol.%) 14 – 8 – 18 12 20 9.5 40.4 30 41 28
ple, in specimen code 1B735(15), 1 stands for steel number, B for the starting microstructure, 735 for ICAT and (15) for MVF.
3.
Results and discussion
3.1.
Microstructures
On heating starting microstructure A to the ICAT, austenite (martensite at room temperature) nucleated at, and grew to consume, the pearlite and then grew into the ferrite. The quenching from different ICAT produced martensitic structure with different MVF restricted to ferrite/ferrite boundaries and an isolated or continuous network of martensite structure along ferrite/ferrite boundaries formed depending on MVFs (Fig. 5(a) and (b)). During intercritical annealing of starting microstructure B, the preferential area for austenite nucleation was the lath interfaces, lath colony boundaries and prior austenite grain boundaries. The austenite volumes obtained on intercritical annealing and then water quenching was then smaller and more finely dispersed (Fig. 6(a) and (b)). This result is in good agreement with literature (Erdogan and Priestner, 1999, 2002; Erdogan, 2002; Hughes, 1986; Jamiru, 1990; Thomas and Koo, 1981; Tomota, 1987). The size and distribution of the martensite obtained on quenching then reflected the scale of the initial microstructure.
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Fig. 5 – Micrograph of 1A735(15) (a) and 2A755(25) (b) specimens. 1A735(15) specimen, intercritically annealed at 735 ◦ C and cooled at 0.3 ◦ C/s to give 17% martensite and 8% new ferrite and 2A755(25) specimen, intercritically annealed at 737 ◦ C and cooled at 5 ◦ C/s to give 25% martensite and 10% new ferrite etched in 2% nital.
Fig. 6 – Micrograph of 2B717(15) (a) and 2B717(25) (b) specimens. 2B717(15) specimen, intercritically annealed at 717 ◦ C and cooled at 0.3 ◦ C/s to give 14% martensite and 12% new ferrite and 2B717(25) specimen, intercritically annealed at 737 ◦ C and cooled at 1400 ◦ C/s to give 26% martensite and 0% new ferrite etched in 2% nital.
Fig. 7 – Quantitative microstructure maps of series A specimens of steel 1 (a) and steel 2 (b) annealed at 735 ◦ C showing effect of cooling rates on the microstructure (old ferrite: ferrite present during intercritically annealing).
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3.2.
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Quantitative microstructure maps
Figs. 7 and 8 illustrate some examples of the effect of cooling rate on the proportion of phases after heat treatment in the form of the microstructure maps for series A and B starting microstructures. Since the martensite and new ferrite transforms from austenite present at ICAT during cooling, the sum of their volume fraction should be equal to the volume fraction of austenite present at the intercritical temperature (although some austenite may be retained in some dual-phase steels). For the present steel, the cooling rates in excess of 1000 ◦ C /s were required to convert all of the austenite into martensite. With decreasing the cooling rates, martensite was increasingly replaced by the new ferrite. At the slower cooling rates, the formation of the new ferrite was followed by the transformation of remaining austenite to martensite and carbide was absent. Some retained austenite was probably present at some cooling rates, as will be discussed later. At any cooling rates, the absence of the ferrite–carbide aggregate formation is an important feature of the transformation characteristics in these steels since the presence of carbide in dual-phase steels has a detrimental effect on ductility. The other important feature of those microstructure maps (Figs. 7 and 8) at the cooling rate of ∼50 ◦ C/s corresponding to the required cooling rate for industrial application of rapid cooling stage IV (Fig. 1), carbide is absent and MVF is over 20%. The absence of carbide formation is unusual at this cooling rate and the critical cooling rates for martensite formation compared with the previous studies is so low. Generally dualphase steels having MVF between 15 and 25% exhibit optimum mechanical properties (Rocha et al., 2005). Some examples of microstructures obtained after cooling from ICAT ranges at different cooling rates are shown in Fig. 9(a) and (b). Cooling the specimen 2B717(15) from 717 ◦ C (Fig. 9(a)) converted most of the austenite to martensite, the
martensite volume fraction being close to ∼15% and the new ferrite content being ∼12%. Cooling the specimen 2B795(15) from the temperature of 795 ◦ C (Fig. 9(b)), resulted in more new ferrite and 17% martensite. In Fig. 9(a) and (b), the chromate etch invented by Lawson et al. (1980) was used to show the new ferrite as the bright, white constituent in the micrographs, 12% in Fig. 9(a) and 39% in Fig. 9(b). The critical cooling rate for martensite formation is quite low when compared with the previous studies; (Lawson et al., 1980, 1981; Eldis, 1979; Matlock et al., 1979; Jeong and Kim, 1985; Marder, 1977; Geib et al., 1980; Mould and Skena, 1977; Rigsbee and Vander Arend, 1977; Huppi et al., 1980). The reason for this may be due to the combination of low carbon and the high Ni contents. As is mentioned before, in the current dual-phase literature, effect of Ni content variation on austenite hardenability has not been reported yet. Mn is well known to have an effect on austenite hardenability and is without any exception present in dual-phase steels with changing amounts in different alloys. However, in the previous studies (Lawson et al., 1980, 1981; Eldis, 1979; Matlock et al., 1979; Jeong and Kim, 1985; Marder, 1977; Geib et al., 1980; Mould and Skena, 1977; Rigsbee and Vander Arend, 1977; Huppi et al., 1980; Erdogan and Priestner, 1999, 2002; Erdogan, 2002; Sarwar and Priestner, 1999), still high amount of ferrite carbide aggregate have been observed to be present even in dual-phase steels which have high Mn contents at slower cooling rates. Therefore, the effect of Ni on austenite hardenability is focused on and underlined here. Figs. 7 and 8 also show that a significant amount of austenite is transformed to martensite even at the cooling rate of 0.01 ◦ C/s (furnace cooling). At the same MVF, much less new ferrite was present after cooling from low intercritical annealing temperatures than after cooling from the highest intercritical annealing temperature (Tables 3 and 4). This is due to the higher carbon concentration in austenite at lower ICAT and consequently higher hardenability 19].
Fig. 8 – Quantitative microstructure maps of series B specimens of steel 1 (a) and steel 2 (b) annealed at 735 ◦ C showing effect of cooling rates on the microstructure (old ferrite: ferrite present during intercritically annealing).
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Fig. 9 – Micrograph of 2B717(15) (a) and 2B795(15) (b) specimens. 2B717(15) specimen, intercritically annealed at 717 ◦ C and cooled at 0.3 ◦ C/s to give 14% martensite and 12% new ferrite and 2B795(15) specimen, intercritically annealed at 795 ◦ C and cooled at 0.54 ◦ C /s to give 17% martensite and 39% new ferrite. New ferrite revealed as bright white areas after etching in hot chromate reagent.
3.3.
Effect of austenite dispersion on its hardenability
The microstructure maps shown in Figs. 7 and 8 were interpreted in the form of austenite → martensite hardenability diagrams (Priestner and Ajmal, 1987), as in Fig. 10(a) and (b). In dual-phase steels, the proportions of the several constituents present after dual-phase heat treatment depend on the cooling rate, the composition of the austenite and the degree of segregation of the substitutional element prior to the cooling from the ICAT. In addition, Priestner (1987) proposed that the dispersion of the austenite in ferrite at the ICAT should have a significant effect on the hardenability. The dispersion of the austenite particles in the ferrite was characterized by the ratio of interfacial area to austenite volume fraction. Priestner showed by geometric analysis that as this ratio increased the yield of martensite at any cooling rate should decrease, because for a given rate of epitaxial ferrite growth, a large interfacial area would result in a larger volume of new ferrite when the temperature fell to the Ms of
the austenite. For a given volume fraction of austenite, a high interfacial area equates to fine austenite particles. Hughes (1986) demonstrated that the starting microstructure defined the morphology of the austenite at the intercritical annealing temperature. His study was on 0.12% C, 1.5% Mn, 0.34% Si steel, which was homogenized at 1150 ◦ C for 20 h and then heat treated to provide starting microstructures of different scales. On intercritical annealing, the different starting microstructures produced a range of austenite dispersions. The finer starting microstructure gave a finer, more dispersed austenite, which needed faster cooling to achieve a constant volume fraction of martensite, nearly the same with Priestner’s proposal. Using the same steel, Jamiru (1990) also investigated the effect of austenite dispersion on its hardenability. From the different starting microstructures a range of austenite dispersions were yielded by intercritical annealing. He also showed that there was a strong reduction in martensite yield as the dispersion parameter increased for all cooling rates. However,
Table 3 – Cooling rates and phase proportions in steel 2 specimens Specimen code 2A717(15) 2A717(25) 2B717(15) 2B717(25) 2A737(15) 2A737(25) 2B737(15) 2B737(25) 2A795(15) 2A795(25) 2B795(15) 2B795(25)
Cooling rate (◦ C/s)
Martensite content (vol.%)
0.3 1400 0.042 1400 0.04 5 0.04 5 0.54 50 0.52 50
14 26 14 26 15.5 25 17 26 17 27 15 26
New ferrite content (vol.%) 12 – 12 – 19.5 10 18 9 39 29 41 30
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j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 8 ( 2 0 0 8 ) 75–84
Fig. 10 – Martensite hardenability of various austenites of steels.
he pointed out that at the slower cooling rates the relation between the dispersion and the martensite content was weaker than the geometric model predicted. The reason for the disagreement between the model and the experimental data was explained as follows; in the model, it was assumed that the cooling rate was constant and that an average ferrite growth velocity applied throughout cooling. In reality this is not the case. Ferrite growth rate is a complex function of the degree of undercooling, carbon diffusion rate and carbon concentration in the austenite at the ␣/␥ interface. The model ignores the fact that small particles of austenite must become enriched in carbon more rapidly than large particles, with respect to the amount of ferrite growth. This suggests that for a given cooling rate, the average ferrite growth rate will be considerably slower for fine dispersions than the coarse ones. Therefore, less ferrite and more martensite will form relative to the geometric theory, and this difference will increase with increasing austenite dispersion. In the present experiments the finer starting microstructure, series B, produced more martensite than the coarser A, after the intercritical heat treatments at the lower temperatures, and particularly at the slower cooling rates, i.e. at the slowest cooling rate of 0.01 ◦ C/s, both MVFs of finer B and coarse A series for steels 1 and 2 are ∼11% (Fig. 8(a)) and ∼7% (Fig. 7(a)) respectively. Under the same cooling condition, finer series B of steel 2 (Fig. 8(b)) were also produced more MVF than coarse series A of steel 2 (Fig. 8(b). After the intercritical heat treatment at the highest temperature, the hardenability of austenite from both series was nearly identical. To explain these results, it is necessary to propose that the carbon-enrichment effect is capable of balancing or overwhelming Priestner’s geometric effect. In the present work, the concentration of the austenite stabilizing elements (1.72–1.89% Mn and 0.67–0.85% Ni) is higher than in Jamiru’s and Erdogan’s (Erdogan, 2003) steel, and would depress the Ms temperature. Also, the carbon content of the present steel is only half that of Jamiru’s steel and almost same as Erdogan’s steel. In fully austenitised steel, the critical cooling rate for 100% martensite is increased by reducing
the carbon content and decreased by increasing the substitutional alloy content. In the present steels, this critical cooling rate after the intercritical heat treatment is approximately 1000 ◦ C/s, which is similar to that for higher carbon, lower alloyed steels investigated previously. However, a substantially larger fraction of martensite is retained to much lower cooling rates in the present steels. This suggests that the alloying elements have a large effect in reducing the ferrite growth rate, despite the lower carbon concentration. However, it will be noted that the ferrite growth must be very sensitive to carbon concentration in the austenite. Therefore, much more new ferrite has been formed in the austenite obtained at high ICAT than the one at low ICAT ones. It is the sensitivity of ferrite growth rate to the increasing carbon concentration in the austenite as the austenite shrinks during the cooling. That negates Priestner’s geometric effect when the materials have been annealed at the high ICAT, but after annealing at the low ICAT, Priestner’s theory has been seen to be approved. When both steels have been cooled from approximately the same temperature of 735 or 737 ◦ C, the required cooling rates to obtain 25% martensite are 1400 ◦ C/s for the steel 1 and only 5 ◦ C/s for the steel 2 (Figs. 7 and 8). The influence of Mn and Ni on the hardenability by depressing the Ms temperature is well known. The amount of Mn and Ni in the steel 2 is higher than the steel 1 and both steels have the same carbon content. In addition to martensite and ferrite, retained austenite may be present in dual-phase steel in amounts varying from 2 to 9% (Rigsbee and Vander Arend, 1977; Yi and Kim, 1983; Jeong and Kim, 1987; Bangaru and Sachdev, 1982; Furukawa et al., 1979; Saleh and Priestner, 2001). The amount of retained austenite is sensitive to the cooling rate, intercritical annealing temperature and carbon content of the steel. Rigsbee and Vander Arend (Rigsbee and Vander Arend, 1977) and Bangaru and Sachdev (Bangaru and Sachdev, 1982) reported that the retained austenite content increased slightly with the increasing ICAT and the carbon content of the steel. It has been observed (Priestner, 1987; Yi and Kim, 1983; Jeong and Kim, 1987) that the amount of the retained austenite also increased as the cooling rate decreased. In the present work the retained
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 8 ( 2 0 0 8 ) 75–84
austenite was found to be barely detectable in some samples cooled slowly.
4.
Conclusions
The following conclusions could be drawn from the above discussions. 1. The critical cooling rate for the martensite formation is considerable low compared with the previous studies. This is probably due to the combination of low carbon and high Ni contents. 2. In the present steels ferrite-carbide aggregates did not form at the regardless of any cooling rates. This result is unusual compared with the previous studies. 3. The fine dual-phase structure produced more martensite than the coarse microstructure after annealing at the low temperature and particular at the slower cooling rates. The reason for this was considered to the result from higher carbon enrichment effect in the fine structure than the coarse ones during the cooling. 4. With considerable low quenching power, the compositions of the present steels are suitable for industrial production of dual-phase steel through continuous annealing line of iron and steel plants, providing a useful quantity of martensite and pearlite, and ferrite-carbide aggregates are absent.
Acknowledgements The authors wish to acknowledge the financial supports of Eregli Iron and Steel Plant, Turkey for this study. The author is also indebted to The Scientific and Technical Research Council of Turkey, High Technology Institute Gebze, Turkey for providing their induction furnace castings and cold rolling facilities.
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