δ phase precipitation and its effect on the mechanical properties of a super duplex stainless steel

δ phase precipitation and its effect on the mechanical properties of a super duplex stainless steel

Materials Science and Engineering, A 174 ( 1994 ) 149-156 149 a Phase precipitation and its effect on the mechanical properties of a super duplex st...

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Materials Science and Engineering, A 174 ( 1994 ) 149-156

149

a Phase precipitation and its effect on the mechanical properties of a super duplex stainless steel J i a n c h u n Li*, T i y a n W u * a n d Y v e s R i q u i e r

General MetallurgyDepartment, FacultePolytechnique de Mons, Mons (Belgium) (Received April 13, 1993; in revised form June 18, 1993)

Abstract A super duplex stainless steel containing nominally (by weight) 25 Cr, 7 Ni, 3.5 Mo and 0.25 N has been investigated for o phase precipitation and its effect on mechanical properties at elevated temperatures. The specimens were solution treated and some of these were annealed between 700 and 1050 °C for times ranging between 1 and 60 min. Selective etching, optical microscopy, X-ray diffraction, bulk extraction and chemical analysis have been used to identify the austenite, ferrite and o phases, and to construct a time-temperature-precipitation diagram for o phase formation. Hot tensile tests and subsequent fractography were performed on specimens containing a wide range of o phase volume fractions. The results indicate the following: the incubation time for o precipitation is about 5 rain at 850-900 °C; the ductility of specimens deformed at about 600 °C is greatly reduced by o phase precipitation; the effects of this phenomenon increase with the o phase volume fraction and exhibit quite a marked temperature dependence; at higher temperatures, e.g. 900 °C, o phase particles in the steel are realigned along the shearing direction during deformation and the final fraction appears ductile

1. Introduction Super duplex stainless steels may be defined as a family of steels having a ferritic-austenitic structure and a pitting resistance equivalent (PRE) greater than 40, where P R E = % C r + 3 . 3 % M o + 16%N. Their low cost and good engineering performance [1, 2] have led to an increasing number of applications, mainly in chemical and offshore industries, where high levels of strength and resistance to corrosion are required. However, a large variety of undesirable phases (carbides, nitrides, intermetallic phases) may appear in these alloys and greatly affect their properties if appropriate manufacturing procedures are not adopted. Among the phases listed above, the o phase is the most important because of its drastic effect on the toughness and corrosion resistance of the material [3, 4]. Therefore, it is very important to understand clearly the o embrittlement mechanisms during production as well as in service. Although a phase formation has received considerable attention in austenitic stainless steels, apart from the work of Maehara and coworkers [5, 6], little infor-

*Present address: Department of Materials, Central-South University of Technology, Changsha, China. 0921-5093/94/$7.00

mation exists about its effect on the mechanical properties in highly alloyed duplex stainless steels. The purpose of this investigation was to examine the o phase morphologies, the kinetics of formation of this phase and its effect on hot ductility in a 25Cr-7Ni-3.5Mo-0.25N super duplex stainless steel after undergoing thermal exposures at temperatures ranging from 600 to 1100 °C.

2. Experimental procedure The material used in the present study was a plate (18 mm thick) produced by hot rolling of a continuously cast slab. The chemical composition of the steel is given in Table 1. A solution treatment temperature of 1100 °C was necessary to dissolve all the o phase in the ferrite, and a standard treatment of 60 min at this temperature, followed by water quenching was chosen. Figure 1 shows the microstructure of the hot-rolled plate after solution treatment. Island-like austenite grains can be seen in the ferrite matrix. Isothermal treatment in the temperature range 700-1050 °C for times up to 100 min resulted in the transformation of ferrite to austenite and the a phase. Hot tensile testing, fractography, hardness measurements, chemical analysis of the o phase and determination of the volume © 1994 - Elsevier Sequoia. All fights reserved

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a Phase precipitation and mechanical properties

TABLE 1. Chemical composition of the steel used (wt.%) C

Si

Mn

P

S

Cu

Ni

Cr

Mo

W

N

B

0.018

0.43

0.7

0.02

0.001

0.57

7.2

25

3.6

0.65

0.25

0.002

servohydraulic testing machine under a constant strain rate of 0.01 s -1. Specimens 10 mm in diameter and 160 mm long were prepared parallel to the initial rolling direction. A heating rate of 100 °C s -1 was obtained by passing an electric current through the specimen. The temperature was measured and controlled by means of a thermocouple spot welded to the center of the specimen's surface. After being fractured, the specimens were quenched by compressed argon to protect the fractured surface against oxidation and to warrant a good scanning electron microscopy (SEM) observation. The microstructure was also examined to reveal the void formation. Fig. 1. Microstructure of hot-rolled plate after solution heat treatment at 1100 °C, followed by water quenching. fractions of 6, 7 and a phases by X-ray diffraction were performed on these specimens. For the microscopic examination, 6, ), and a phases were colored gray, white and black, respectively, by electrolytic etching in K O H solution. Vickers microhardness measurements with a load of 50 g were employed to determine the hardness of each phase. The a phase fraction was determined by X-ray diffraction as follows. The ratios of the diffraction intensities of the (331 ), (411 ), (212) and (410) planes to those of the "reference sample" annealed at 800 °C for 100 min were averaged to eliminate the effect of texture. In the "reference sample", austenite and a phases--not ferrite--were detected. The a phase fraction in this "reference" was determined by a bulk extraction technique. 10% HCI in methanol solution was used as the electrolytic extraction medium. No phases other than the a phase was detected in the extracted powder. The 6/7 ratio in the plate samples was determined using the (211) and (200) reflections for ferrite and (311), (220) and (200) for austenite. Co K a radiation was used. Analysis of the alloying elements in each phase was carried out to reveal the partitioning of the elements during annealing. The elements were analysed by means of energy-dispersive X-ray spectroscopy (EDXS) on the 6 and 7 grains. The composition of the a phase was determined chemically from the extraction residues. Isothermal, hot tensile tests over the temperature range 600-1100 °C were carried out on a Gleeble

3. R e s u l t s and d i s c u s s i o n

3.1. Precipitation during isothermal annealing The results obtained from X-ray diffraction and optical microscopy have been used to establish the fractional changes of the 6, ), and o phases with the isothermal annealing time. An example of the data relative to samples annealed at 800 °C is shown in Fig. 2. When annealing at temperatures below 1000 °C, the first precipitate detectable in X-ray diffraction, bulk extraction and optical micrographs consists of austenite. After longer times, the reaction 6--'~+ a

follows the austenite precipitation. When the a phase volume fraction increases, that of ferrite decreases and that of austenite increases. The results obtained from X-ray diffraction have been used to construct the time-temperature-precipitation diagram shown in Fig. 3. In addition to the a phase, the starting point of austenite precipitation (about 3% increasing in austenite volume fraction, as detected by X-ray diffraction) is also presented. The curves of o phase precipitation are of C-type, with a nose at 850-900 °C. At these temperatures, the a phase appears within 5 min and ferrite decomposes completely after 30 min. Although the time-temperature-precipitation diagram is similar to those of austenitic-ferritic duplex stainless steels [3, 7], the precipitation in this steel is faster than those in other cases. This is probably as a result of the higher molybdenum and tungsten contents, which are found to favor a phase formation much more strongly than does chromium [4].

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3.2. Morphology and chemical composition of the phase Figures 4(a)-4(c) show the morphology of the o phase at different stages of precipitation. In an austenitic-ferritic duplex stainless steel, o phase particles normally nucleate at the ferrite-austenite interfaces and grow into the adjacent ferrite grains, developing a cellular structure consisting of the o phase and new austenite. In the former austenite, as far as the optical microscopic observation is concerned, only traces of the o phase can be identified at the grain boundaries after complete ferrite decomposition. The EDXS results for the ferrite and austenite of the specimen after solution treatment are shown in Table 2, and the results of chemical analysis of the o phase precipitates extracted after isothermal annealing at 800 °C for different times are shown in Table 3. From these data listed in Tables 2 and 3, the chromium and molybdenum enrichments in the o phase are remarkable. The variation in chemical composition of the o phase precipitates indicates that the molybdenum content is a function of time. As the molybdenum content in the o phase increases, the chromium content decreases. 3.3. Effect of o phase precipitation on mechanical properties Figure 5 shows the effects of the o phase and temperature on the tensile properties for specimens containing no o phase (as-solution-treated) and 16% o phase (annealed for 20 min at 800 °C). The variations in strength and ductility with temperature are significant. In both cases, the strength decreased with increasing temperature. However, the characteristics of the reduction in area vs. the temperature are quite different

in the two cases examined. For a specimen containing 16% o phase, the ductility decreases progressively with increasing temperature over the entire temperature range investigated. In contrast, for the as-solutiontreated sample, the ductility decreases with increasing temperature before rising again, so that a minimum appears on the curve of ductility vs. temperature. The microstructures in longitudinal cross-section ch)se to the fractured surface after testing at 1100 and 600 °C are illustrated in Figs. 6(a) and 6(b), respectively, for the specimens without any o phase. Although the reductions in area are reasonably high (95% and 70%), the grain size of austenite in these two cases is quite different. When deforming at 600 °C, austenite islands were elongated extensively along the tensile axis, while, at 1100 °C, the austenite islands underwent less deformation before fracture occurred. This can be seen by comparing Figs. 6(a) and 6(b) with Fig. 1. The Vickers hardness of the y and 6 phases at room temperature in the deformed zone are HV 230 (7), 260 (6) and HV 320 (y), 260 (6), respectively, before and after tensile testing at 1000 °C. These values indicate that ferrite could be fully dynamically recovered during testing, while austenite partially retains its work hardening. In the temperature range 600-1000 °C, the selfdiffusion rate in the ferrite is about 100 times that in the austenite [8]. The remarkable difference in recovery kinetics could be responsible for this argument. Figure 6(b) shows that, at about 600 °C, both the ferrite and austenite were plastically deformed to approximately the same degree. When the temperature rises, ferrite can recover more easily than austenite during the deformation process. The difference in deformation behavior between the 7 and 6 phases increases and the ductility of the duplex stainless steel decreases with temperature. Despite this, the ductility is improved at even higher temperatures, e.g. 1100 °C, where ferrite dominates the deformation process.

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The effect of the o phase on the mechanical properties is significant at lower temperatures, e.g. 600 °C; however, this effect quite strongly depends on temperature, as can be seen in Fig. 5. At higher temperatures, e.g. 900 °C, the ductility and ultimate tensile strength (UTS) for both solution-treated and annealed specimens are very near. The reduction in area against temperature curve for the annealed specimen shows no

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TABLE 3. Phase composition (wt.%), determined by chemical analysis, for specimens annealed at 800 °C Annealing time (min)

Si

Ni

Cr

Mo

Cu

8 12 20 100

0.48 0.50 0.53 0.58

3.50 3.51 3.73 3.72

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5.70 7.02 7.86 8.35

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Mo

Cu

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a Phase precipitation and mechanical properties

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turning point. T h e ductility increases monotonically with temperature. In this case, the a phase in the annealed steel is the dominant factor in the ductility. T h e difference in deformation behavior between the ferrite and austenite phases, which caused a minimum in the ductility in the as-solution-treated specimen, was suppressed here by the hard and brittle a phase. To evaluate the effect of the a phase content on the mechanical properties, the reduction in area and UTS at two temperatures are plotted as a function of the o phase fraction in Fig. 7. At 600 °C, the reduction in area decreases from 70% to as low as 5% when the fraction of a phase raises from 0% to 26%. However, at 900 °C, this variation is not significant. Only a slight decrease in ductility is shown when the fraction of o phase increases from 0% to 26%. 3.4. F r a c t o g r a p h y

A typical example of a fractured surface in the solution-treated condition tested at 600 °C is shown in Fig. 8. Transgranular ductile dimples, which originated

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Fig. 7. Variation of tensile properties at 600 and 900 °C with volume fraction of a phase.

Fig. 8 Fracture surface of specimen after solution treatment at 1100 ~C. followed by water quenching, and deformation at 600 °C.

from the decohesion of the inclusion-matrix interfaces, exhibit high ductility. Indeed, deep dimples with fine structure lines are clearly shown in Fig. 8. T h e size and distribution of these lines correspond to the deformed

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cr Phase precipitation and mechanical properties

substructures, such as slip bands or subgrains, in the austenite as well as in the ferrite. The longitudinal section of this sample, which appears to be of excellent quality, is shown in Fig. 6(b). In contrast, the fracture surface of the same kind of specimen exhibits relatively poor ductility when ruptured at 900 °C (Fig. 9(a)). The longitudinal section of this fracture is also shown in Fig. 9(b). In this case, a large number of voids were formed around the )'-6 interfaces close to the fractured surface. These voids were linked by shear fracture of the ferrite, with cracks propagated in directions along the tensile axis as well as at about 45 ° to it. Austenite islands are surrounded by the more ductile ferrite at 900 °C. The decohesion of austenite from the ferrite matrix has occurred and can be clearly recognized, as indicated by the arrows in Fig. 9(b). Only the ferrite phase is responsible for the final ductile fracture. Less deformation substructures

are observed on the dimples in Fig. 9(a) (unlike the fracture at 600 °C, where both phases took part in the final fracture and more plastic deforamation has taken place before failure). The decline in ductility with increasing temperature in the range 6 0 0 - 8 0 0 °C can also be explained in terms of this fracture mechanism. Figures 10(a) and 10(b) present the fracture surface and longitudinal section of a specimen annealed at 800 °C for 20 min (16% o phase) and tested at 600 °C. Very low elongation has been observed. The appearance of the fracture surface was granular and no ductile dimples could be observed. Cleavage cracks were linked by thin shearing fracture, where austenite was plastically deformed; however, ductile austenite could not compensate for the brittleness of the a phase. Many "river" patterns are visible on the cleavage facets, as shown by the arrows in Fig. 10(a). In the longitudinal section (Fig. 10(b)), transverse micro-

Fig. 9. (a) Fracture surface and (b) corresponding microstructure in the longitudinal cross-section of specimen after solution treatment at 1100 °C and deformation at 900 °C.

Fig. 10. (a) Fracture surface and (b) corresponding microstructure in the longitudinal cross-section of specimen after solution treatment at 1100 °C, annealing at 800 °C for 20 min and deformation at 600 °C.

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cracks across the o-austenite cellular structures are formed by brittle fracture. However, a specimen with the same structure as that shown in Fig. 10(b) exhibited ductile fracture when deformed at 900 °C, as can be seen in Fig. 1 l(a). The longitudinal section of this specimen close to the fracture is shown in Fig. 1 l(b). In this case, although voids formed by the decohesion of ~-austenite interfaces, the orientation of the o phase had been changed during the deformation process. By comparing Fig. 1 l(b) with Fig. 4(b), it is clear that the o phase particles had realigned in the direction about 45 ° to the tensile axis after deformation. This contrasts with Fig. 4(b), where the o phase bands are perpendicular to the d - 7 interface. The realignment of the a phase demonstrates that quite an amount of ductility was preserved in the

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former ferrite region, where large amounts of ~ phase particles existed. On the fracture surface, fine dimples corresponding to the shearing process are displayed in Fig. l l(a). Furthermore, when a specimen with the same structure as that shown in Fig. 10(b) was deformed at 1000 °C, plastic shearing predominates in the fracture surface, which declines in a direction about 45 ° to the tensile axis, indicating excellent ductility (Fig. 12(a)). The corresponding microstructure (Fig. 12(b)) showed that no a phase was present. Nevertheless, a very fine secondary austenite phase can be observed in the former ferrite regions. There are two possible reactions responsible for this phenomenon. Firstly, the o phase is dissolved at 1000 °C but the secondary austenite formed by the eutectoid reaction d - - ' o + 7 remains partially. Alternatively, ferrite, which is

Fig. 11. (a) Fracture surface and (b) corresponding microstructure in the longitudinal cross-section of specimen after solution treatment at 1100 °C, annealing at 800 °C for 20 min and deformation at 900 °C.

Fig. 12. (a) Fracture surface and (b) corresponding microstructure in the longitudinal cross-sections of solution treatment at 1100 °C. followed by water quenching, annealing at 800 °C for 20 min and deformation at 1000 °C.

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thermodynamically unstable at 1000 °C, may transform to austenite. This process would be accelerated by plastic deformation during testing. At this point, ferrite was hardened slightly by the dispersion of secondary austenite. This results in a higher UTS of the annealed specimen than that of a specimen without previous o phase precipitation, as can be seen in Fig. 5. Granular roughening on the fracture surface corresponds to this dispersion in some places, as indicated by an arrow in Fig. 12(a). An example of the microhardness measurements is shown in Fig. 12(b), where the sizes of the Vickers hardness indentations within the ferrite and austenite demonstrate clearly the difference between these two phases. After hot deformation at 1000 °C, austenite is much harder than ferrite, although it is well known that, in the solution-treated condition, these values are reversed.

4. Conclusions (1) The o phase has C-curve transformation kinetics with a minimum transformation time of the order of 5 min at 850-900 °C. (2) After nucleation at the austenite-ferrite interface, the o phase grows into the ferrite in a cellular morphology. As the o phase grows, its molybdenum content increases and its chromium content decreases. (3) The ductility of the steel in the solution-treated condition presents a minimum value at 800-900 °C.

(4) The o phase embrittlement depends on the a phase content even more strongly than on the deformation temperature. (5) At temperatures below 600 °C, the ductility is greatly reduced by o phase precipitation. (6) During deformation at higher temperatures, e.g. 900 °C, the o phase particles are realigned in the shearing direction and the final fracture exhibits a ductile morphology with a large reduction in area. (7) The effects of o phase precipitation on the mechanical properties disappear at 1000 °C.

Acknowledgment The authors wish to thank Fabrique de Fer de Charleroi for supplying the alloy used in this study.

References 1 H. D. Solomon and T. M. Devine, in R. A. Luda (ed.), Duplex Stainless Steel, American Society for Metals, Metals Park, OH, 1983, pp. 553-572. 2 T.A. DeBold,Z Met. (March 1989) 12-15. 3 Y. Maehara, M. Koike, N. Fujino and T. Kunitake, Trans. ISIJ, 23(1983) 240-246. 4 L. A. Norstrom, S. Pettersson and S. Nordin, Z. Werkstofftech., 23(1983) 220-234. 5 Y. Maehara and Y. Ohmori, MetalL Trans. A, 18 (1987) 663-672. 6 Y. Maehara, Y. Ohmori and T. Kunitake, Metal. Technol., lO (1983) 296-303. 7 E G. E. Beetge and E E A. Robinson, Met. Mater., (September 1973) 408-413. 8 J. Barcik and B. Brzycka,Met. Sci., 17(1983) 256-260.