Effect of Heat Treatment on the Mechanical Properties of ... - CiteSeerX

Journal of ASTM International, January 2005, Vol. 2, No. 1 Paper ID JAI13037 Available online at www.astm.org

Marcelo Martins1 and Luiz Carlos Casteletti2

Effect of Heat Treatment on the Mechanical Properties of ASTM A 890 Gr6A Super Duplex Stainless Steel ABSTRACT: Duplex stainless steels were initially developed for the manufacture of thermomechanically formed products. However, the increasing demand for casting components of these materials has led to the application of a widely developed technology for this forming process. After undergoing a solution annealing heat treatment, these materials become thermodynamically metastable systems, since the concentration of solute atoms in solid solution is so high that they become saturated, causing them to seek a lower free energy state when exposed to different temperatures. These systems reach a more stable thermodynamic condition by the precipitation of various intermetallic phases, depending on the temperatures to which they are exposed. Thermal energy serves as a catalyst to “overcome” the energy barrier that separates the metastable and stable phases. The objective of this work was to determine the influence of several heat treatment temperatures on the microstructure and mechanical properties of an ASTM A 890/A 890M Gr6A super duplex stainless steel. The increase in hardness and the decrease in impact toughness of these materials in impact tests were found to be directly correlated with the increase in sigma phase concentration in their microstructure, which tended to precipitate into ferrite/austenite interfaces. When the sigma phase was completely dissolved by the heat treatment, the material’s hardness was determined by the volumetric concentration of ferrite and austenite in the microstructure, and the energy absorbed in the impact test reached approximately 220 J at room temperature. KEYWORDS: sigma phase, microstructure, super duplex stainless steel

Introduction Duplex stainless steels are widely used in “offshore” platforms, since they offer good mechanical properties and excellent pitting corrosion resistance. Despite the greatly varying concentrations of salt, hydrogen sulfide (H2S), carbon dioxide (CO2), sulfide ions such as (HS-) and (S-2), and oxygen (O2) contents in seawater, duplex stainless steels are still a good option for such applications [1]. Duplex stainless steels include a particular category called super duplex stainless steels (SDSS) whose microstructure is composed of ferrite and austenite, each present in a volumetric concentration of about 50 ± 5 %. SDSSs present a Pitting Resistance Equivalent number (PREN) higher than 40 [2], a property which is theoretically expressed by Eq 1 [3]. PREN = {%Cr +[(3,3).( %Mo)] + 16(%N)}

(1)

Duplex and super duplex stainless steels are Fe-Cr-Ni-Mo-N alloys containing up to 0,3wt.% of nitrogen in atomic form and displaying a similar corrosion resistance and a superior mechanical performance to that of copper alloys [2]. Manuscript received 18 August 2004; accepted for publication 15 October 2004; published January 2005. 1 M.Sc., Industrial Manager, SULZER Brasil S/A - Fundinox Division, Av. Eng. João F. Gimenez Molina, 905 Jundiaí - SP - Brazil, CEP 13213-080 - Cx Postal 2114. Professor of Centro Universitário Salesiano (Unisal). 2 Ph.D., Associate Professor, Department of Materials, Aeronautics and Automotive Engineering, University of São Paulo’s São Carlos School of Engineering.

Copyright © 2005 by ASTM International, 100 Barr Harbor Drive, PO Box C700, West Conshohocken, PA 19428-2959.

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The main properties of duplex and super duplex stainless steels represent a combination of the best characteristics of their two phases, with the austenite contributing to their impact toughness and the stronger ferrite improving their tensile strength and weldability. These steels offer advantages over austenitic stainless steels, displaying a higher stress corrosion resistance and a greater intercrystalline corrosion resistance, as well as superior mechanical properties. Their mechanical properties depend on the ferrite content in the microstructure, whose normal concentration varies from 60 % to 40 % in volume. Higher ferrite (δ) contents increase the mechanical strength but favor the precipitation of sigma phase during solidification cooling through a eutectoid reaction of the type: δ ⇌σ+γ

(2)

This reaction is unbalanced toward the right, i.e., in the direction of sigma phase formation, as the concentration of ferrite-stabilizing elements, particularly chromium and molybdenum, increases [2]. It is practically impossible to prevent the formation of sigma phase during solidification cooling, because the compositional ranges favor its precipitation; however, the volumetric fraction of that phase can be minimized by adjusting the chemical composition and cooling rate during the solidification process. In addition to the sigma phase, complex chromium and molybdenum carbides, R phase, G phase, ε phase (in copper containing steels), and chromium nitrides, among others, can also appear in these materials’ microstructure. Due to these steels’ strong tendency for secondary phase precipitation, it is essential to study the temperatures at which these materials become sensitive. The “incubation” time for precipitation of the σ phase is 5 min between 850°C and 900°C. The material’s ductility is strongly reduced by σ phase precipitation. The effect of this phenomenon increases as the volumetric fraction of σ phase increases. The σ phase formation mechanism occurs through controlled nucleation and subsequent, relatively fast growth. The σ phase is the most exhaustively studied of all the intermetallic phases that appear in these steels’ microstructure, for, in addition to appearing in a higher proportion, it drastically affects the material’s impact toughness and corrosion resistance. Therefore, we have attempted to investigate the influence of the several heat treatment temperatures on the phase proportions, hardness, and energy absorbed in impact tests. Experimental Procedure The test specimens were cast in vacuum induction degassing (VID) furnaces operating at the electric supply system’s frequency (60 Hz) and at a maximum power of 400 KW. The first stage consisted of the casting design of the test specimens, having a 25 mm diameter and 300 mm length, using AUTOCAD 2000 equipped with MECHANICAL DESKTOP software, followed by a simulation of the solidification process using a specific software program called SOLSTAR. The chemical composition of the material was analyzed via optical emission spectrometry. Based on the chemical composition, the values of Crequivalent and Niequivalent were calculated according to the following expressions [4]:

MARTINS AND CASTELETTI ON HEAT TREATMENT OF STEEL

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Creq.= (%)Cr + [(1,5). (%)Si] + [(1,4). (%)Mo] + (%)Nb – 4,99

(3)

Nieq. = (%)Ni + [(30). (%)C] + [(0,5). (%)Mn] + [26(%)N - 0,02)] + 2,77

(4)

The solution annealing heat treatment recommended for this material was performed in an electric furnace with a heating capacity of 1300ºC. The solution annealing heat treatment at 1160°C was carried out to completely dissolve all the precipitates in the ferritic matrix, forming an unsaturated solid solution. After this solution annealing heat treatment, followed by water quenching, the samples were aged at various temperatures ranging from 520–1180°C, increasing the treatment temperatures in 20°C steps, to check the influence of these temperatures on the phase concentrations and on the energy absorbed in the impact test. All the samples were kept for 2 h at the aging temperatures. For the quantitative analysis, the samples were etched with three different etchants: Behara II, Murakami, and a 50–50 % solution of HCl and HNO3, using full immersion. The percentage of each phase in the microstructure was determined by the grid point counting method. The method adopted here involved the use of a grid containing 100 points and 200 times magnification, based on the ASTM E 562 standard. A Stereoscan 440– LEO scanning electron microscope (SEM) was also used, and the images by secondary electrons were taken in digital form. This device was equipped with a 3PC Microspec-type EDS (Energy Dispersive Spectrometer) detector. After being heat treated, the samples were subjected to a hardness test to ascertain the influence of the quantities of phases on this property of the material. The samples were also subjected to an impact test at room temperature. Results and Discussion The chemical compositions of the samples studied here are shown in Table 1. After the chemical analysis, the Creq/Nieq ratio obtained from Eqs 3 and 4 resulted in 1.71, corresponding to a theoretical ferrite percentage of approximately 48.2 % and, hence, to 51.8 % of austenite, according to the ASTM A 800/A 800M standard [4]. TABLE 1—Concentration of chemical elements (wt %) indicated by spectrometry. C (%) Cr(%) Mo(%) Ni(%) Si(%) Mn(%) Cu(%) W(%) 0,016 25,69 3,80 7,18 0,74 0,52 0,716 0,736 Nb(%) Ti(%) Al(%) V(%) Zr(%) Co(%) Sn(%) Pb(%) 0,014 0,005 0,016 0,049 0,065 0,055 0,0069 0,0018

optical emission N(%) 0,22 S(%) 0,008

P(%) 0,027 Fe(%) Rem.

The delta ferrite content, indicated by quantitative metallography using the grid point technique, was 43 % in volume, representing a downward variation of 10.8 %, according to the value forecasted by ASTM A 800. The delta ferrite content in the microstructure depends on the heat treatment temperature, since the volumetric concentration of ferrite increases as this variable rises. The volumetric fraction of delta ferrite in the microstructure also depends on the correlation between the ferrite and austenite stabilizing elements and on the heat treatment temperature. The figures below depict the materials’ microstructures under different heat treatment conditions after solution annealing at 1160°C followed by water quenching.

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The sample solution, annealed at 1160°C, quenched, and treated at 520°C for 2 h, displays a microstructure composed of only two phases: the ferritic matrix and the austenite precipitated in the form of “islands,” as illustrated in Fig. 1. No intermetallic phases are precipitated at the δ/γ interfaces or at the ferrite grain boundaries. The sample annealed at 1160°C and treated at 820°C for 2 h shows the presence of austenite, sigma phase, and ferrite dispersed in the sigma phase and austenite mixture, as depicted in Fig. 2.

γ

δ

FIG. 1—Scanning electron micrograph of the sample treated at 520°C for 2 h.

γ

σ+δ+γ

FIG. 2—Scanning electron micrograph of the sample treated at 820°C for 2 h. Ferrite is difficult to identify because at that temperature, its volumetric concentration is very low, i.e., about 3.0 %. However, as can be noted, the austenite appears on a plane below that of the sigma phase plane because the latter is harder than the austenite and during mechanical


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polishing, it stands out in high relief since it becomes less worn. The austenite obtained from the eutectoid reaction is mixed with the ferrite, but it is difficult to identify each of these phases at this level of magnification. In the sample treated at 920°C (Fig. 3), the visible sigma phase is clearly outlined and in relief for the aforementioned reason. At that temperature, this phase clearly precipitates in a continuous network. This morphology favors the propagation of cracks at the interfaces between the phases. The treatment at 940°C (Fig. 4) produced two types of morphologies in the sigma phase: an elongated one in relief, with well-defined boundaries, and the second, lacy one encompassing austenite and delta ferrite, which provided chromium and molybdenum for its precipitation.

γ

σ+γ


γ

δ+σ+γ σ


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The major difficulty involved in the manufacture of these steels is to obtain the maximum PREN value without impairing the microstructural balance between ferrite and austenite. When super duplex stainless steels are exposed to temperatures below the solution annealing temperature, the metastable thermodynamic balance is disturbed, causing the system (material) to seek a more stable thermodynamic state through the precipitation of intermetallic phases, carbonic phases, and a microstructural imbalance between ferrite and austenite. These intermetallic phases include the sigma phase, which is very rich in chromium and molybdenum. Figure 5 shows the correlation between the proportions of the phases and the diverse heat treatment temperatures. Austenite

Ferrite

Sigma Phase

Phase Volumetric Concentrations (%)

70 60 50 40 30 20 10 0 500

600

700

800

900

1000

1100

1200

Heat Treatm ent Tem peratures (°C)

FIG. 5—Proportions of the phases versus heat treatment temperature. As can be seen, the sigma phase begins to precipitate at 720°C, and its concentration rises sharply up to 800°C, when it reaches about 45 % in volume. The concentration peak of this phase occurs close to 880°C, with a volumetric concentration of 50 %. Between 880°C and 900°C, the percentage of sigma phase drops by 10 % in volume. From 900–1000°C, the volumetric concentration of the sigma phase gradually decreases, displaying a sharp drop between 1000°C and 1020°C. The sigma phase gradually disappears from the material’s microstructure at temperatures above 1060°C. The curve indicating the variation in volumetric concentration of ferrite displays an inverse behavior to that representing the concentration of sigma phase. An increasing concentration of sigma phase causes the concentration of ferrite to decrease because the elements forming the intermetallic phase (Cr and Mo) are rejected from the ferrite, where they are more soluble. This causes a depletion of these elements in the ferrite and a considerable decrease in the Crequiv/Niequiv ratio, favoring its transformation into secondary austenite. Thus, the eutectoid reaction: δ ⇌ σ + γ2 is unbalanced toward the right. When the heat treatment temperature rises from 1000°C to 1020°C, the volumetric concentration of ferrite increases abruptly from 5 % to 38 %, then stabilizes at around 46 % between 1140°C and 1180°C.


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A practical way to detect the presence of sigma phase in the microstructure is to measure the material’s hardness, as can be seen by a comparison of Figs. 5 and 6. Figure 6 shows that the steel’s hardness begins to increase effectively from 720°C on, displaying a peak in the order of 400 Brinell at 800°C, which remains high up to treatments at 840°C, thereafter diminishing as the temperature rises. Figure 6 also reveals that the hardness peak at 800°C does not coincide with the sigma phase concentration peak, which occurs at 880°C. A possible reason for this is the fact that, with the eutectoid decomposition of the delta ferrite into sigma phase and secondary austenite, the ferritic phase disappears completely, in detriment to the two new phases. Thus, as a result of the significant increase in austenitic phase, the hardness decreases considerably. At temperatures exceeding 1060°C, the material’s hardness tends to stabilize between 250 and 260 Brinell, which represents hardness typical of a super duplex stainless steel [7]. 450

Brinell Hardness (HB)

400

350

300

250

200 500

600

700

800

900

1000

1100

1200

Heat Treatment Temperatures (°C)

FIG. 6—Hardness of the material as a function of heat treatment temperatures. Figure 7 illustrates the material’s hardness as a function of the volumetric percentage of sigma phase. As the volume of sigma phase increases, the material’s hardness increases with an approximately parabolic tendency. This intermetallic precipitate is hard, leading to a significant increase in this property from the macroscopic standpoint. Figure 08 shows the impact toughness values measured through a Charpy test conducted at room temperature on a V-notched test specimen. Note that the energy absorbed in the test drops to 10% of the original value, with a sigma phase volume of only 3% in the microstructure. Figure 09 illustrates the behavior of impact toughness as a function of the heat treatment temperature. As can be seen, the energy absorbed in the test drops abruptly from 540oC on, reaches a value of 20J at 640oC, remains at extremely low values (below 10J) up to 1040oC, then “leaps” rapidly to the level of 200J from 1060oC on, remaining at around that value at higher temperatures. Moreover, even without the presence of the sigma phase, i.e., with a volumetric percentage equal to zero, there is a visible drop in absorbed energy during the impact test. This effect was not detected either in the hardness versus heat treatment temperature curve through an increase in this property, or in the micrographs obtained by optical microscopy.


BRINELL HARDNESS VS SIGMA PHASE CONCENTRATION

BRINELL HARDNESS

450

400

350

300

250

200 0

10

20

30

40

50

60

VOLUMETRIC CONCENTRATION OF SIGMA PHASE (%)

FIG. 7—Correlation between Brinell hardness and sigma phase volume.

250 ABSORBED ENERGY (J)

8

200 150 100 50 0 0

10

20

30

40

50

60

VOLUMETRIC CONCENTRATION OF SIGMA PHASE (%) FIG. 8—Energy absorbed in Charpy Test versus sigma phase volumetric percentage.


9

250

Absorbed Energy (J)

200

150

100

50

0 500

600

700

800

900

1000

1100

1200

Heat Treatment Temperature (°C)

FIG. 9—Absorbed energy versus heat treatment temperature. According to Fig. 5, the sigma phase begins precipitating from 720°C on, so the drop in absorbed energy at temperatures below 720°C (see Fig. 9) cannot be attributed to the precipitation of this intermetallic phase. The analyses of the fractured surfaces of samples treated at 520°C, 980°C, and 1160°C are shown in Figs. 10–15. Figures 10 and 11 show the appearance of the fractures in samples treated at 520°C. Note the fibrous aspect typical of highly tough materials, which indicates that the heat treatment temperature did not modify the mechanical behavior in terms of impact toughness. This finding is also confirmed by an analysis of the absorbed energy versus temperature curve (Fig. 9). This fracture surface displays included particles originating from the deoxidization process, as well as voids around the particles. Figure 12 shows the peaks revealed by a qualitative analysis of the particle obtained by EDS detector. As can be seen, this is an inclusionary particle of a complex silicon, zirconium, and aluminum oxide originating from the deoxidization process. Figure 13 depicts the appearance of the fracture in the sample treated at 980°C for 2 h. At this temperature, the material displays a brittle behavior due to the presence of a large amount (approximately 31 %) of sigma phase in the microstructure. The appearance of the fracture differs totally from the previous case, lacking the presence of fibrous structures. Note the appearance of an intergranular crack likely due to the precipitation of sigma phase in that region. The appearance of the fracture differs slightly from the typical pattern for brittle materials since, in this case, there is a high volumetric fraction of precipitated austenite (58 %), in detriment to the reduction of the volumetric fraction of ferrite. This high content of a ductile and tough phase alters the fracture appearance pattern for brittle materials. Figures 14 and 15 show aspects of the fractures in samples treated at 1160°C. Note the presence of a fibrous structure typical of ductile material at that treatment temperature, indicating the absence of intermetallic phases negatively influencing the impact toughness. The presence of tough phases in the microstructure indicates that part of the energy required to break the test specimen is absorbed, plastically straining the material and hindering crack propagation.

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FIG. 10—Appearance of the fracture in the sample treated at 520°C for 2 h.

FIG. 11—Appearance of the fracture in the sample treated at 520°C for 2 h.

MARTINS AND CASTELETTI ON HEAT TREATMENT OF STEEL 11

FIG. 12—Qualitative analysis of the particle shown in Fig. 11.

FIG. 13—Appearance of the fracture of the sample treated at 980°C for 2 h.


FIG. 14—Appearance of the fracture in the sample treated at 1160°C.

FIG. 15—Appearance of the fracture in the sample treated at 1160°C. Conclusions The sigma phase begins precipitating at heat treatment temperatures in the order of 720°C, dissolving completely from solution annealing above 1060°C, followed by water quenching. The sigma phase precipitates preferentially at the ferrite/austenite interface, growing toward ferrite, which supplies elements such as chromium and molybdenum for its stabilization. At the heat treatment temperature of 880°C, the volumetric concentration of sigma phase reaches the maximum value of 50 % in volume.

MARTINS AND CASTELETTI ON HEAT TREATMENT OF STEEL 13

The volumetric fraction of ferrite drops to very low values, gradually disappearing at heat treatment temperatures in the 840°C to 880°C interval. The material’s hardness begins to increase significantly from 720°C up, reaching its highest values at treatment temperatures of 800–820°C. This peak does not coincide exactly with the peak of maximum volumetric concentration of sigma phase, which occurs at 880°C. At high treatment temperatures exceeding 1060°C, the material’s hardness is affected only by the ferrite and austenite phases. The material’s hardness is strongly influenced by the volumetric fraction of sigma phase in the microstructure. The energy absorbed in the impact test is strongly reduced by the presence of sigma phase in the microstructure. A mere 3 % in volume of sigma phase in the microstructure reduces the absorbed energy to 10 % of its original value, i.e., when this intermetallic phase is absent from the microstructure. After heat treatments at high temperatures (above 1060°C), the energy absorbed in the impact test returns to the material’s original levels. References [1] Duplex Stainless Steels ’91, Charles, J. and Bernhardsson, S., Eds., Beaune, Les Ulis, France, Les Éditions of Physique, Vol. 1, 1991, pp. 3–48. [2] Weber, J., “Materials for Seawater Pumps and Related Systems,” Sulzer Brothers Limited, Winterthur, Switzerland, pp. 1–11. [3] ASTM Standard A 890/A 890M–91, “Ferrous Castings; Ferralloys,” Annual Book of ASTM Standards, Vol. 01.02, ASTM International, West Conshohocken, PA, 1999, pp. 556–569. [4] ASTM Standard A 800/A 800M–91, “Ferrous Castings; Ferralloys,” Annual Book of ASTM Standards, Vol. 01.02, ASTM International, West Conshohocken, PA, 1999, pp. 458–463. [5] ASTM Standard A 781/A 781M–91, “Ferrous Castings; Ferralloys,” Annual Book of ASTM Standards, Vol. 01.02, ASTM International, West Conshohocken, PA, 1999, pp. 440–443. [6] ASTM Standard A 370–97a, “Ferrous Castings; Ferralloys,” Annual Book of ASTM Standards, Vol. 01.02, ASTM International, West Conshohocken, PA, 2001, pp. 148–200. [7] Weber, J. and Schlapfer, H. W., “Austenitic-Ferritic Duplex Steels,” Sulzer Brothers Limited, Winterthur, Switzerland, pp. 1–10. [8] Maehara, Y., Ohmori, Y., Murayama, J., Fujino, N., and Kunitake, T., “Effects of Alloying Elements on σ Phase Precipitation in δ-γ Duplex Stainless Steels,” Metal Science, Vol. 17, November 1983, pp. 541–547. [9] Li, J., Wu, T., Riquier, Y., “σ Phase Precipitation and Its Effect on the Mechanical Proprieties of the Super Duplex Stainless Steel,” Materials Science and Engineering, 1994, Vol. 174A, pp. 149–156. [10] Anson, D. R., Pomfret, R. J.,and Hendry, T., “Prediction of the Solubility of Nitrogen in Molten Duplex Stainless Steel,” ISIJ International, Vol. 36, No. 7, 1996, pp. 750–758. [11] Goldstein, J. I., et al., “Scanning Electron Microscopy and X-Ray Microanalysis,” 2nd ed., New York, Plenum Press, Cap. 3, 1992, pp. 69–147. [12] Rossitti, S. M., “Effect of Niobium in the Microstructure and in Mechanical Properties of the Molted Super Duplex Stainless Steel SEW 410 W. Nr. 1.4517,” Thesis, Universidade de São Paulo, Interunidade, EESC-IFSC-IQSC, 2000, pp. 49–53. [13] Weiss, B. and Stickler, R., “Phase Instabilities at High Temperature Exposure of 316 Austenitic Stainless Steel,” Metallurgical Transaction, Vol. 3, April 1972, pp. 851–864.


[14] Pohl, M., “The Ferrite/Austenite Ratio of Duplex Stainless Steels,” Z. Metallkd., Vol. 86, No. 2, May 1995, pp. 97–101. [15] Siewert, T. T., McCowan, C. N., and Olson, D. L., “Ferrite Number Prediction to 100 FN in Stainless Steel Weld Metal,” Welding Research Supplement, December 1988, pp. 289-s– 297-s.