Abnormal creep behavior of ferritic Fe24Cr4Al stainless steel

Pergamon

Scripta Metallurgicaet Materialia,Vol. 30, No. 11, pp. 1397-1402, 1994 Copyright© 1994ElsevierScienceLtd Printed in the USA. All fightsreserved 0956-716X/94 $6.00 + 00

ABNORMAL CREEP BEHAVIOR OF FERRITIC FE-24CR-4AL STAINLESS STEEL S.C. Tjong and J.S. Zhang Department of Physics and Materials Science City Polytechnic of Hong Kong, Tat Chee Avenue, Kowloon, Hong Kong

(Received December 30, 1993) (Revised February 15, 1994) Introduction The creep behavior of solid solution alloys has been divided into two groups based on the value of the stress exponent and other creep characteristics: the alloy class (Class I) and the pure metal class (Class II) (13). It is believed that dislocation climb is the rate-controlling step during deformation in pure metals whereas viscous glide motion of dislocations dragging the solute atmosphere is the rate-controlling process in class I materials. The stress exponent is about equal to 3 for Class I materials and it is approximately equal to 5 for Class II alloys. Furthermore, it has been suggested that Class I behavior is favoured when the misfit in size between the solute and solvent atoms is large (2). There are several potential advantages in using ferritic Fe-Cr-A1 alloys in high temperature applications, i.e. lower raw materials cost, superior oxidation resistance, and lower coefficient of thermal expansion than the Ni-base superalloys (4). In view of these conditions, ferritic Fe-Cr-A1 alloys find some applications in high temperature oxidation environment such as heating elements, nuclear reactors, petroleum refineries and automotive exhaust system (5). The cyclic deformation behavior of this alloy has been investigated previously (6, 7). It was found that the labyrinth structures that are typical substructures developed in fatigued f.c.c. metals also tend to form in this ferritic alloy. The presence of persistent slip bands (PSBs) in the ferritic alloy would drastically reduce its fatigue life. This investigation has been principally aimed to establish the parameters necessary for predicting the high temperature creep behavior of Fe-24Cr-4Al alloy. The addition of Al to Fe-24Cr alloy for enhancement of the oxidation resistance can lead to abnormal creep behavior. Experimental Ferritic Fe-24Cr-4A1 alloy used in this investigation has a composition of 23.85 wt. % Cr, 3.89 wt. % Ai, 0.009 wt. %C and the balance iron. The alloy was melted in a vacuum induction furnace and then hot forged into plates of 25 mm thickness. Square bars with a cross-section of - 25 x 25 mm were cut from these plates. The bars were solution treated at 1323 K for 1 h followed by water quenching. The average grain size of the material after the heat treatment is -630/~m. Creep specimens were machined from the bars. The specimens used had a gauge length of 50 nun with a diameter of 8 ram. Tensile creep tests under constant load were carried out in the stress range of 40 - 100 MPa at 873-923 K. Creep strain was monitored continuously using an extensometer incorporating linear variable displacement transducer with the extensometer knife edges attached to the specimen grips. The temperature of the specimen was measured with Pt/Pt-Rh thermocouples. The temperature was kept constant to within + 1 K. Specimens for transmission electron microscopy were sectioned from the specimens crept to a steady state followed by rapid cooling at constant load. Strain transient dip tests [8] were performed at 873 K for an initial stress of 100 and 80 MPa, 1397

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respectively. In this test, the sample is crept under a constant initial stress into steady state. The applied stress is suddenly reduced and held constant at the new reduced stress level. The subsequent creep rate is measured. Results In all test conditions of stresses and temperatures, an inverse primary creep curves are observed. Fig. 1 shows typical creep data of the Fe-24Cr-4A1 alloy at 873 K for various stresses in the form of the logarithm of true creep rate vs. the logarithm of time. The amount of instantaneous strain upon load application is small, thus the initial creep rate of the alloy is very low. The rate then increases continuously with time until it attains a steady state. Class I alloys generally exhibit such a primary creep behavior (1). However, it should be noted that in typical b.c.c. Class I alloy, a steady state is reached after experiencing a relatively large strain, e.g., 0.2 - 0.3 for Fe-Mo alloy (10). However, the strain for attaining a steady state is very small in the present alloy and it is less than 0.02 in most cases. Figure 2 shows the steady creep rate of the alloy as a function of applied stress at 873 and 923 K, respectively. It can be seen in this figure that the creep data obey the creep power law with a stress exponent (n) of 5.57 at 923 K which increases to a value of 6.0 at 873 K. Fig. 3 shows a temperature dependence of the steady creep rate at 60 and 70 MPa, respectively. The activation energy for creep (Qc) calculated from the slopes of the straight lines is 310 - 330 kJ/mol. This value is greater than the reported value of the activation energy for AI diffusion in Fe (Qs = 246 kJ/mol) and is close to the reported activation energy for lattice diffusion of Fe (Qi = 290 kJ/mol) (9). Fig. 4 shows the relationship between the steady creep rate normalized by the activation energy for creep, Qc, and applied stress normalized by Young's modulus, E, measured at test temperatures. Figure 5 summarizes the results of the strain transient dip tests. It is apparent in this figure that the creep rate obtained immediately following stress reductions is higher than the steady state creep rate at the reduced stress level. According to the literature, it appears that Class I alloys tend to exhibit this type of transient behavior following stress reductions (10). It has been reported that the creep rate of Class I alloys is controlled by dislocation glide dragging its solute atmosphere and is directly proportional to the dislocation density (11), i.e. = 9bY

(i)

where p is the mobile dislocation density, b is the Burgers vector and v is the average dislocation velocity which is determined by the rate at which the solute atoms can move along with the dislocations. In the present work, the fast creep rate following stress reduction can be explained in terms of the present of high density of dislocations in the specimen deformed at an initial high stress prior to stress reduction. The subsequent decrease in creep rate is associated with a gradual decrease in the dislocation density to a value expected under the reduced stress condition. If the dislocation structure is assumed to remain constant when the stress suddenly changed, the constant structure creep rate can be obtained as a function of stress by determining the creep rate immediately after stress reduction from an initial stress to various stresses (Fig. 5). The results are shown in Fig. 6 in which a fourth power stress dependence of the constant structure creep rate is observed. These results together with the sixth power stress dependence of steady state creep rate lead to a conclusion that the structure changes as a second power of the applied stress. This is reasonable if the present alloy is assumed to be a Class I alloy. As there is no cell or subgrain structure formed in Class I alloy during the steady state creep deformation, thus the dominant structural parameter characterizing the creep structure is the dislocation density which would increase as the square of the applied stress. Figure 7 consists of TEM micrographs of the Fe-24Cr-4AI alloy crept to the steady state at 873 K under the applied stress of 80 and 60 MPa, respectively. It is apparent in these micrographs that the dislocations are slightly curved and they exhibit planar distribution. There is no tendency to form the cell or subgrain structure. The dislocation density in the specimen crept at 80 MPa is higher than that crept at 60 MPa (Fig. 7(b)). All these dislocation configurations observed are typical dislocation structures of Class I alloys. In Fig. 7(c), fine particles precipitate extensively along grain boundaries. The structure and compositions of the precipitates have been identified by the electron diffraction and EDS analysis as M23C6.

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Discussion Experimental results mentioned above show that the present alloy exhibit abnormal creep behavior. Primary creep and transient creep behavior and the dislocation structure exhibit typical Class I characteristics, but the stress and temperature dependence of the steady state creep rate is similar to that of Class II alloys. The Class I creep behavior is usually observed in solid solution alloys having a large size difference between the solute and solvent atoms, which leads to a strong interaction between the solute atoms and the dislocation [2]. Thus the present alloy can be considered as a Class I alloy because the difference in the atomic size between the solute A1 and solvent Fe is comparable to that of typical Class I alloys such as AI-Mg and Fe-Mo alloys. The discussion will be focused on the reason why the stress and temperature dependence of the steady state creep rate deviates from the characteristics of Class I alloy. One reason we might consider is the grain boundary precipitation. It has been shown that in an anstenitic stainless steel (Class II alloy) grain boundary precipitation increases the stress exponent of steady state creep rate from 5 to 7 (12). This has been attributed to the presence of the grain boundary particles which strongly obstruct dislocation annihilationat the grain boundaries, leading to an increase of the dislocation density near the grain boundaries. In Class I alloys, however, this effect is expected to be small, if any, because the creep rate is controlled by viscous glide of dislocations dragging their solute atmosphere but not by dislocation annihilation. In fact, we did not observe any dislocation pile up at the grain boundaries or an increase in the dislocation density near the grain boundaries of the present alloy (Fig. 7(c)). Another fact we might consider is the effect of the addition of third alloying element Cr in the specimen. All the Class I alloys investigated to date are binary alloys, and there are no data available concerning the effect of third alloying element. As the atomic size and Young's modulus of the third element Cr are almost identical to those of solvent Fe (9), its effect on the interaction between the solute A1 and dislocations would be negligible. A number of models (11, 14) on the creep rate of Class I alloys have been proposed by several workers. Although different assumptions were made on the dislocation glide processes, these models simpl,y yielded the same results, i.e., dislocation density p a 03 and viscous glide velocity v ct tr, thus ~ = pbv tx o~, accounting for the commonly observed n value of 3. Our results of the constant structure creep (Fig. 6) and the dislocation structure observations (Fig. 7) have shown that the dislocation density obeys the second power law of the applied stress, i.e., pcx o2. This indicates that the dislocation velocity seems not to be proportional to o. Friedel (14) has proposed a model for the case in which the solute atmosphere around the dislocation is not saturated and the creep rate equation can be expressed as ~s

= 2p -b-b D s S i n h

°b21s

(2)

kT

where p is the mobile dislocation density, 13, is the solute diffusion coefficient, b is the Burgers vector and ks is the mean distance between solute atoms dragged along a dislocation. ob2ts In this equation the dislocation glide velocity is proportional to sinh ( " - ~ ) but not o and only when, aleX, < < kT, this equation is reduced to ~ c~ bpe c~ 03. According to this model, for the present test conditions of relatively low temperature ( < 0.5Tm) and high stress ( > 10~/~), the stress exponent of 6, is expected to be larger than 3.

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It can also be seen from Eq. 2 that the apparent activation energy for creep, Qc, depends on the stress and Qc is greater than the solute diffusion activation energy, Qs. The value of Qc obtained in this work is also stress dependent and Qc is greater than Qs. This is in good agreement with the prediction of the Friedel model. The Friedel model is valid for viscous glide of dislocation which is not saturated by solute, i.e., for relatively low solute concentrations around the dislocations. Because of the absence of data for the interaction energy between the solute and dislocation in ternary alloys, we cannot confirm at present whether the solute atmosphere is saturated or not and further investigation is needed to elucidate this problem.

Conclusion Fe-24Cr-4AI ferritic stainless steel exhibits abnormal creep behavior. Primary creep, transient creep after stress reduction and the dislocation structure exhibit Class I creep characteristics but the stress and temperature dependence of the steady state creep rate exhibit Class II creep characteristics. If the solute atmosphere around dislocations is assumed to be not saturated, the abnormal creep behavior can be explained according to the Friedel model for Class I alloys. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.

O.D. Sherby and P.M. Burke, Prog. Mater. Sci. 13, 325 (1968) W.R. Cannon and O.D. Sherby, Metall Trans. 1, 1030 (1970) F.A. Mohamed and T.G. Langdon, Acta Metall. 22, 779 (1974) F.G. Wilson, B.R. Knott and C.C. Desforces, Metall Trans. 9A, 275 (1978) S.D. Sastry, P.K. Rohatgi, K.B. Abraham, and Y.V.R.K. Prasay, Scripta Metall. 13,817 (1979) S.C. Tjong, L.T. Wu and N.J. Ho, Mater. Sci. Eng. 100, 79 (1988) S.C. Tjong, I.C. Hsieh and N.J. Ho, Z. Metallkde 79, 189 (1988) C.N. Ahlquist and W.D. Nix, Scripta Metall. 3,679 (1969) Metals Data Handbook, Japan Institute of Metals, Tokyo (1984), in Japanese R.W. Evans and B. Wilshire, Creep of Metals and Alloys, p. 122, The Institute of Metals, London (1985) J. Weertman, Acta MetaU. 25, 1393 (1977) J.S. Zhang, P.E. Li and J.Z. Jin, Acta Metall. Mater. 39, 3063 (1991) S. Takeuchi and A.S. Argon, Acta Metall. 24, 883 (1976) J. Friedel, Dislocations, Addison Wesley, Reading, Mass. (1964), p 239, 409. Acknowledgements

This work was supported by UPGC Competitive Earmarked Grant (Grant No. 904054, CPHK 227/92E). J.S. Zhang was on leave from Dalian University of Technology, China.

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