- Email: [email protected]
TEM studies of dislocation substructure in 316 austenitic stainless steel strained after annealing in various environments
Materials Science and Engineering A249 (1998) 91 – 96
TEM studies of dislocation substructure in 316 austenitic stainless steel strained after annealing in various environments Witold Zielin´ski, Ayad A. Abduluyahed, Krzysztof J. Kurzydłowski * Department of Materials Science and Engineering, Warsaw Uni6ersity of Technology, Narbutta 85, 02 -524 Warsaw, Poland Received 27 October 1997; received in revised form 10 February 1998
Abstract Results of transmission electron microscopy investigations of specimens annealed in either air or vacuum at 600°C and strained to a total elongation of 0.5% at room temperature are presented. A macroscopically nonuniform spatial distribution of dislocations has been revealed. The dislocation density appears to be larger in the vicinity of the surface of the deformed specimens which indicates a higher work hardening rate near the scale – matrix interface. The dislocation density is also nonuniform along a plane parallel to the interface. In this case, however, the variations in the dislocation density seem to be random. The higher values of the flow stress determined in the vacuum-annealed specimens may be attributed to the higher strength of the oxide scale developed in such conditions of annealing (low pressure of the oxygen). © 1998 Elsevier Science S.A. All rights reserved. Keywords: Austenitic stainless steel; Environmental effect; Oxide scale; Dislocation structure
1. Introduction It has been known for a long time that crystal plasticity depends on the state of the surface of specimens. This phenomenon has been a subject of extensive studies in the conditions of monotonic as well as cyclic deformation [1,2]. In the case of austenitic stainless steel (SS) the plastic deformation in the temperature range of serrated flow is not uniform. This behaviour, known as the Portevin– LeChatelier effect (PLC), is usually explained in terms of dynamic strain ageing, which arises due to the diffusion of solute atoms to the moving dislocations [3 –6]. It has been also attributed to the precipitation of chromium carbides which may lead to a local variation of concentration of chromium and carbon [7]. Mechanical properties of metals in general depend not only on the test temperature but also on the test environment [8,9]. On the other hand, studies of PLC effect usually are carried out in air. Recently, Ab* Corresponding author. Tel.: +48 22 6607420; fax: + 48 22 484875; e-mail: [email protected] 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(98)00578-4
duluyahed et al. [10,11] and Kurzydłowski [12] have reported on the effect of test environment on the flow stress in 316 austenitic SS at the temperatures of dynamic and static ageing. The results demonstrated a difference in the resistance to plastic deformation between the specimens of 316 austenitic SS tensile strained in: (a) air and vacuum at 600°C, and (b) at room temperature after annealing in air and vacuum at 600°C. The increased values of the flow stress observed systematically after annealing/straining in vacuum were discussed in relation to the thickness and chemistry of the oxide layers formed on the specimens and their influence on the properties of the material. The discussion of the environmental effect on the mechanical properties of austenitic SS presented in Refs. [10,11] was concentrated on the characterisation of the specimen surface and, in particular, on the role of the oxide layers. The present investigation was aimed at describing the dislocation substructures in the base material (substrate). Observations were carried out on the material in the near-surface zone and in the centre of specimens annealed in various test environments and strained at room temperature. The results give a better
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
92
Table 1 Chemical composition of the steel examined (wt.%) C
Cr
Ni
Mo
Mn
Si
Cu
Ti
Al
P
S
Fe
0.042
17.00
15.90
2.65
2.03
0.19
0.17
0.01
0.03
0.023
0.005
Balance
insight into the processes responsible for effect of test environment on the plastic flow in austenitic SS.
2. Experimental details The study was carried out on a 316-type stainless steel delivered in the form of cylindrical rods (see also Refs. [10–12]). The chemical composition of the material is given in Table 1. Flat tensile specimens with gauge section dimensions of 10×5 × 2 mm were annealed at 900°C in air for 1 h. As a result, a singlephase microstructure with a grain size of approximately 10 mm was obtained. As-recrystallized specimens were subsequently annealed at 600°C either in vacuum (1.33 × 10 − 4 Pa) or in air (atmospheric pressure) and strained in tension at room temperature to a total strain of 0.5% at a strain rate of 2× 10 − 4 s − 1. All tensile tests were carried out on the same universal testing machine manufactured by Instron. The machine was calibrated before each test and a number of specimens were used to ensure reproducibility of the data. (The stress – strain curves typical of the specimens annealed at 600°C in vacuum and air
Fig. 1. Room temperature stress–strain curves for 316 stainless steel specimens recrystallized at 900°C and subsequently annealed at 600°C in vacuum and air.
as well as those recrystallized at 900°C are shown in Fig. 1). Room temperature straining was applied in order to activate the potential dislocation sources present in the as-annealed specimens. Dislocation substructures in the strained materials were studied using a transmission electron microscope under an accelerating voltage of 100 kV. Each thin foil was examined under several tilt angles with respect to the electron beam. Thin foils were prepared by a double jet electropolishing or dimpling and ion milling, or by combination of these two methods. Several cylinders 3 mm in diameter were cut out in the direction perpendicular to the surface of the flat samples using spark machining. The cylinders were sliced with a precise wire saw in order to obtain specimens for transmission electron microscopy (TEM) observations representative of the material at a specified distance from the free surface. Precise measurements of the slice thickness allowed estimation of the distance from the specimen surface for each TEM thin foil. Assuming that the double jet electropolishing removes the same amount of the material from both sides of the slice, it can be concluded that TEM observations were carried out at the following distances from the outer surface of the specimens: (a) 909 50, (b) 360 950 and (c) 1000 950 mm (the latter corresponds to the half-thickness of the bulk samples). In the investigations of scale–matrix interface region, TEM thin foils had to be prepared in a special way. The slices were ion milled (the scale side up) in argon under the following conditions: 1.7 kV, an angle of 5°. The ion milling was carried out until only a few ‘patches’ of the scale were left at the centre of the polished area. After ion milling, the surface was protected with varnish and the samples were polished electrochemically from the matrix side. The specimens for TEM observations of the crosssection of the scale were prepared in a standard way using Gatan G-1 epoxy to stick two pieces of the material together against their oxidised surfaces. A spark machine was used to cut out a cylinder with the axis oriented along the epoxy layer. The cylinder was glued into a tube and then sliced with a precise wire saw to obtain disks 0.1 mm thick and 3 mm in diameter. The TEM specimens were prepared using a Gatan Dimple Grinder 656 and a Precision Ion Polishing System PIPS 691.
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
93
Fig. 2. TEM images of the microstructures of 316 stainless steel recrystallized at 900°C in air for 1 h, annealed at 600°C in vacuum (a–d) and in air (e – h), and subsequently strained up to 0.5% at room temperature. Left column specifies the distance from the surface.
3. Results and discussion TEM observations concentrated attention on: (a) the dislocation arrangement as a function of the distance from the scale–matrix interface, (b) the dislocation arrangement in a plane parallel to the scale–matrix interface, (c) the substructure of the oxide scale.
3.1. Dislocation arrangement as a function of the distance from the free surface Fig. 2 shows TEM photographs typical of the samples annealed in vacuum and air and subsequently strained up to 0.5% at room temperature. One can see that the microstructures of the samples annealed in vacuum (see Fig. 2a – d) and in air (see Fig. 2e –h) look remarkably similar. In both cases, there is, however, a great difference in the dislocation density between the
region close to the surface and at the core of the specimens. (The microstructure of the regions close to the surface is exemplified by the images representative for the material at 360 mm, while the microstructure of the core is at 1000 mm from the surface.) Details of the microstructures in the vicinity of the scale–matrix interface in air- and vacuum-annealed samples are shown in Fig. 2a,e. The density of dislocations is high and the dislocations are mostly uniformly distributed, which is due to the multi-slip system deformation. Microstructures observed at a distance of 90 mm from the surface are shown in Fig. 2b,f. Although the density of dislocations is still high, certain local fluctuations caused by other microstructural features, e.g. grain boundaries, can be observed. Fig. 2c,g shows the microstructures formed at a distance of 360 mm where the dislocation density is uniform. The operating individual slip systems indicate an early stage of the
94
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
Table 2 Estimated dislocation density in the prestrained specimens annealed in vacuum and air Distance from the surface (mm)
Interface 90 9 5 360 95 1000 (core)
Dislocation density (×1012 m−2) Vacuum
Air
42 92 14 91 99 1 29 0.5
53 9 2 12 91 7 91 B1
evolution of the dislocation microstructure. The last two micrographs, Fig. 2d and Fig. 2h, present a number of individual dislocation pile-ups, which suggest early stages of the plastic deformation in the core of the specimens. The dislocation density was estimated as a function of the test environment and the distance from the specimen surface using basic methods of quantitative microscopy (e.g. Ref. [13]). It should be noted that TEM observations of thin foils result usually in underestimation of dislocation density. Consequently, the true numerical values of dislocation density might be higher than reported here. However, the error committed in the case of the series of specimens examined in the present study is approximately constant, and relative values of the density at the selected observation fields presented in Table 2 are reliable. An interesting result obtained in the present investigations is a large dislocation density near the surface. It should also be noticed, that the dislocation density decreases rapidly with increasing distance from the oxide–substrate interface. This indicates that, in early stages of plastic deformation of polycrystalline specimens examined, the surface affects the plastic strain distribution. The observed dislocation distribution confirms that, in the region near the oxide – substrate interface (e.g. in the vicinity of the outer surface), the effect of strain hardening is stronger. The results also
prove, that the scale–matrix interface features either high activity dislocation sources or their increased density. As for the origin of this considerable variation in dislocation density, the present study does not give a direct answer and further research is in progress. At the present moment this phenomenon seems to be linked both to the stress relaxation under applied load and to the relaxation of internal stresses due to the differences in the properties, including volume mismatch, of oxide films and the substrate.
3.2. Dislocation arrangement along the oxide–substrate interface Further details of TEM observations in a plane parallel to the scale–matrix interface are shown in Fig. 3. Although, in all the samples, whether annealed in vacuum or in air, the deformation is macroscopically uniform, areas with highly nonuniform dislocation density are observed. This implies that local stress concentrators are present in the scale or in the region near the scale–matrix interface. The dislocation density changes drastically from a very low level, of 1012 m − 2, up to an upper limit of 1015 –1016 m − 2 along a short distance of several microns (see Fig. 3). At this stage of the investigations no quantitative analysis of the density of these areas was attempted. It has however been found that their density in the thin foils prepared from air-annealed samples is significantly higher than that observed in vacuum-annealed samples.
3.3. Microstructure of the oxide scale In order to explain how the scale influences the plastic deformation of specimens an attempt was made to characterise the microstructures of the oxides developed as a result of annealing at 600°C in air and vacuum. Specimens for TEM observations of the crosssection of the scale developed in vacuum are relatively easy to prepare, whereas the specimens of the scale
Fig. 3. An example of the nonuniform dislocation density in a plane parallel to the scale – matrix interface in the samples of 316 stainless steel annealed in vacuum and in air.
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
95
At this stage of investigations one may conclude that the tough and perfectly dense continuous outer layer of the scale formed on the samples annealed in vacuum improves their mechanical properties compared to the scale on the samples annealed in air, which exhibit a poor adhesion to the metal substrate.
4. Conclusions
Fig. 4. A cross-section TEM image of the scale on 316 stainless steel annealed in vacuum.
produced in air appeared to be extremely difficult because of its poor adhesion to the matrix. Fig. 4 shows a montage of TEM images of the cross-section of a vacuum-deposited scale. In the upper part of the figure an amorphous epoxy layer (used for gluing) is visible. Below the epoxy, there is a layer of uniform material with a submicron grain size of about 0.2 mm. The thickness of the layer is of the order of 1 mm. Beneath this layer, there is a thinner layer with an even smaller grain size. While the interface between these two layers is rather flat, the interface between the second layer and the metal substrate is profoundly rough. It is interesting to note that the layer next to the substrate seems to be porous, contrary to the outer layer, which is perfectly dense. The difficulty in preparing TEM specimens of the scale on the air-annealed samples results from its poor adhesion to the substrate. Fig. 5 shows the substrate (lower half of the picture) and the layer of scale with rather complicated structure (upper part of the picture). A part of the scale, where polycrystalline material is present, is similar in character to the upper layer of the scale formed on the samples annealed in vacuum.
The results obtained in the present work lead to the following conclusions. (1) The distribution of the dislocation density determined in the samples annealed in vacuum and in air suggests that the surface features affect the local plastic deformation and, thus, the mechanical properties of austenitic stainless steel. (2) The high dislocation density near the scale–matrix interface of strained specimens indicates the presence of the active dislocation sources in the vicinity of the free surface. (3) The oxide scale developed on the specimens after annealing at 600°C in vacuum consists of a layer of perfectly dense, submicron polycrystalline material. (4) The oxide scale developed on the specimens after annealing at 600°C in air is characterised by poor adhesion to the substrate. (5) The higher value of the flow stress determined in the vacuum-annealed specimens may perhaps be attributed to the high strength and good adhesion of the oxide scale.
Acknowledgements This work was supported by the State Committee for Scientific Research (KBN) Grant No. 7 T08A 008 12.
Fig. 5. A TEM image of a cross-section of the scale formed on 316 stainless steel annealed in air.
96
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
References [1] R.M. Latanision, J.T. Fourie (Eds.), Surface Effects in Crystal Plasticity, Noordhoff, Leyden, 1977. [2] Viewpoint set No. 18 on: Surface Effects in Cyclic Deformation and Fatigue, Scripta Metall. Mater. 26 (1992) 1499. [3] K.G. Samuel, S.L. Mannan, Rodriguez, Acta Metall. 8 (1988) 2323. [4] A.W. Sleeswyk, Acta Metall. 6 (1958) 599. [5] C.F. Jenkins, G.V. Smith, Trans. Metall. Soc. AIME 245 (1969) 2149. [6] I.S. Kim, M.C. Chaturvedi, Metals Sci. 13 (1979) 691. [7] E.O. Hall, Yield Point Phenomena in Metals and Alloys, Macmillan (Eds.), UK, 1970.
.
[8] T. Nambu, J. Fukumori, M. Morinaga, Y. Matsumoto, T. Sakaki, Scripta Metall. Mater. 32 (1995) 407. [9] M. Morinaga, Y. Murata, M. Furui, T. Wada, Scripta Metall. Mater. 6 (1997) 699. [10] A.A. Abduluyahed, K. Rozniatowski, K.J. Kurzydłowski, Scripta Metall. Mater. 33 (1995) 1489. [11] A.A. Abduluyahed, K. Sikorski, K. Rozniatowski, K.J. Kurzydłowski, in: F.W. Poulsen, N. Bonanos, S. Linderoth, M. Mogensen, B. Zachav-Christiansen (Eds.), Proceedings of 17th Riso Int. Symp. on Mater. Sci., Roskilde, Denmark, 1996, pp. 155 – 160. [12] K.J. Kurzydłowski, Mater. Sci. Eng. A234 – 236 (1997) 1083. [13] K.J. Kurzydłowski, B. Ralph, Quantitative Description of Microstructure of Materials, CRC Press, Boca Raton, FL, 1995.
TEM studies of dislocation substructure in 316 austenitic stainless steel strained after annealing in various environments Witold Zielin´ski, Ayad A. Abduluyahed, Krzysztof J. Kurzydłowski * Department of Materials Science and Engineering, Warsaw Uni6ersity of Technology, Narbutta 85, 02 -524 Warsaw, Poland Received 27 October 1997; received in revised form 10 February 1998
Abstract Results of transmission electron microscopy investigations of specimens annealed in either air or vacuum at 600°C and strained to a total elongation of 0.5% at room temperature are presented. A macroscopically nonuniform spatial distribution of dislocations has been revealed. The dislocation density appears to be larger in the vicinity of the surface of the deformed specimens which indicates a higher work hardening rate near the scale – matrix interface. The dislocation density is also nonuniform along a plane parallel to the interface. In this case, however, the variations in the dislocation density seem to be random. The higher values of the flow stress determined in the vacuum-annealed specimens may be attributed to the higher strength of the oxide scale developed in such conditions of annealing (low pressure of the oxygen). © 1998 Elsevier Science S.A. All rights reserved. Keywords: Austenitic stainless steel; Environmental effect; Oxide scale; Dislocation structure
1. Introduction It has been known for a long time that crystal plasticity depends on the state of the surface of specimens. This phenomenon has been a subject of extensive studies in the conditions of monotonic as well as cyclic deformation [1,2]. In the case of austenitic stainless steel (SS) the plastic deformation in the temperature range of serrated flow is not uniform. This behaviour, known as the Portevin– LeChatelier effect (PLC), is usually explained in terms of dynamic strain ageing, which arises due to the diffusion of solute atoms to the moving dislocations [3 –6]. It has been also attributed to the precipitation of chromium carbides which may lead to a local variation of concentration of chromium and carbon [7]. Mechanical properties of metals in general depend not only on the test temperature but also on the test environment [8,9]. On the other hand, studies of PLC effect usually are carried out in air. Recently, Ab* Corresponding author. Tel.: +48 22 6607420; fax: + 48 22 484875; e-mail: [email protected] 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(98)00578-4
duluyahed et al. [10,11] and Kurzydłowski [12] have reported on the effect of test environment on the flow stress in 316 austenitic SS at the temperatures of dynamic and static ageing. The results demonstrated a difference in the resistance to plastic deformation between the specimens of 316 austenitic SS tensile strained in: (a) air and vacuum at 600°C, and (b) at room temperature after annealing in air and vacuum at 600°C. The increased values of the flow stress observed systematically after annealing/straining in vacuum were discussed in relation to the thickness and chemistry of the oxide layers formed on the specimens and their influence on the properties of the material. The discussion of the environmental effect on the mechanical properties of austenitic SS presented in Refs. [10,11] was concentrated on the characterisation of the specimen surface and, in particular, on the role of the oxide layers. The present investigation was aimed at describing the dislocation substructures in the base material (substrate). Observations were carried out on the material in the near-surface zone and in the centre of specimens annealed in various test environments and strained at room temperature. The results give a better
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
92
Table 1 Chemical composition of the steel examined (wt.%) C
Cr
Ni
Mo
Mn
Si
Cu
Ti
Al
P
S
Fe
0.042
17.00
15.90
2.65
2.03
0.19
0.17
0.01
0.03
0.023
0.005
Balance
insight into the processes responsible for effect of test environment on the plastic flow in austenitic SS.
2. Experimental details The study was carried out on a 316-type stainless steel delivered in the form of cylindrical rods (see also Refs. [10–12]). The chemical composition of the material is given in Table 1. Flat tensile specimens with gauge section dimensions of 10×5 × 2 mm were annealed at 900°C in air for 1 h. As a result, a singlephase microstructure with a grain size of approximately 10 mm was obtained. As-recrystallized specimens were subsequently annealed at 600°C either in vacuum (1.33 × 10 − 4 Pa) or in air (atmospheric pressure) and strained in tension at room temperature to a total strain of 0.5% at a strain rate of 2× 10 − 4 s − 1. All tensile tests were carried out on the same universal testing machine manufactured by Instron. The machine was calibrated before each test and a number of specimens were used to ensure reproducibility of the data. (The stress – strain curves typical of the specimens annealed at 600°C in vacuum and air
Fig. 1. Room temperature stress–strain curves for 316 stainless steel specimens recrystallized at 900°C and subsequently annealed at 600°C in vacuum and air.
as well as those recrystallized at 900°C are shown in Fig. 1). Room temperature straining was applied in order to activate the potential dislocation sources present in the as-annealed specimens. Dislocation substructures in the strained materials were studied using a transmission electron microscope under an accelerating voltage of 100 kV. Each thin foil was examined under several tilt angles with respect to the electron beam. Thin foils were prepared by a double jet electropolishing or dimpling and ion milling, or by combination of these two methods. Several cylinders 3 mm in diameter were cut out in the direction perpendicular to the surface of the flat samples using spark machining. The cylinders were sliced with a precise wire saw in order to obtain specimens for transmission electron microscopy (TEM) observations representative of the material at a specified distance from the free surface. Precise measurements of the slice thickness allowed estimation of the distance from the specimen surface for each TEM thin foil. Assuming that the double jet electropolishing removes the same amount of the material from both sides of the slice, it can be concluded that TEM observations were carried out at the following distances from the outer surface of the specimens: (a) 909 50, (b) 360 950 and (c) 1000 950 mm (the latter corresponds to the half-thickness of the bulk samples). In the investigations of scale–matrix interface region, TEM thin foils had to be prepared in a special way. The slices were ion milled (the scale side up) in argon under the following conditions: 1.7 kV, an angle of 5°. The ion milling was carried out until only a few ‘patches’ of the scale were left at the centre of the polished area. After ion milling, the surface was protected with varnish and the samples were polished electrochemically from the matrix side. The specimens for TEM observations of the crosssection of the scale were prepared in a standard way using Gatan G-1 epoxy to stick two pieces of the material together against their oxidised surfaces. A spark machine was used to cut out a cylinder with the axis oriented along the epoxy layer. The cylinder was glued into a tube and then sliced with a precise wire saw to obtain disks 0.1 mm thick and 3 mm in diameter. The TEM specimens were prepared using a Gatan Dimple Grinder 656 and a Precision Ion Polishing System PIPS 691.
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
93
Fig. 2. TEM images of the microstructures of 316 stainless steel recrystallized at 900°C in air for 1 h, annealed at 600°C in vacuum (a–d) and in air (e – h), and subsequently strained up to 0.5% at room temperature. Left column specifies the distance from the surface.
3. Results and discussion TEM observations concentrated attention on: (a) the dislocation arrangement as a function of the distance from the scale–matrix interface, (b) the dislocation arrangement in a plane parallel to the scale–matrix interface, (c) the substructure of the oxide scale.
3.1. Dislocation arrangement as a function of the distance from the free surface Fig. 2 shows TEM photographs typical of the samples annealed in vacuum and air and subsequently strained up to 0.5% at room temperature. One can see that the microstructures of the samples annealed in vacuum (see Fig. 2a – d) and in air (see Fig. 2e –h) look remarkably similar. In both cases, there is, however, a great difference in the dislocation density between the
region close to the surface and at the core of the specimens. (The microstructure of the regions close to the surface is exemplified by the images representative for the material at 360 mm, while the microstructure of the core is at 1000 mm from the surface.) Details of the microstructures in the vicinity of the scale–matrix interface in air- and vacuum-annealed samples are shown in Fig. 2a,e. The density of dislocations is high and the dislocations are mostly uniformly distributed, which is due to the multi-slip system deformation. Microstructures observed at a distance of 90 mm from the surface are shown in Fig. 2b,f. Although the density of dislocations is still high, certain local fluctuations caused by other microstructural features, e.g. grain boundaries, can be observed. Fig. 2c,g shows the microstructures formed at a distance of 360 mm where the dislocation density is uniform. The operating individual slip systems indicate an early stage of the
94
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
Table 2 Estimated dislocation density in the prestrained specimens annealed in vacuum and air Distance from the surface (mm)
Interface 90 9 5 360 95 1000 (core)
Dislocation density (×1012 m−2) Vacuum
Air
42 92 14 91 99 1 29 0.5
53 9 2 12 91 7 91 B1
evolution of the dislocation microstructure. The last two micrographs, Fig. 2d and Fig. 2h, present a number of individual dislocation pile-ups, which suggest early stages of the plastic deformation in the core of the specimens. The dislocation density was estimated as a function of the test environment and the distance from the specimen surface using basic methods of quantitative microscopy (e.g. Ref. [13]). It should be noted that TEM observations of thin foils result usually in underestimation of dislocation density. Consequently, the true numerical values of dislocation density might be higher than reported here. However, the error committed in the case of the series of specimens examined in the present study is approximately constant, and relative values of the density at the selected observation fields presented in Table 2 are reliable. An interesting result obtained in the present investigations is a large dislocation density near the surface. It should also be noticed, that the dislocation density decreases rapidly with increasing distance from the oxide–substrate interface. This indicates that, in early stages of plastic deformation of polycrystalline specimens examined, the surface affects the plastic strain distribution. The observed dislocation distribution confirms that, in the region near the oxide – substrate interface (e.g. in the vicinity of the outer surface), the effect of strain hardening is stronger. The results also
prove, that the scale–matrix interface features either high activity dislocation sources or their increased density. As for the origin of this considerable variation in dislocation density, the present study does not give a direct answer and further research is in progress. At the present moment this phenomenon seems to be linked both to the stress relaxation under applied load and to the relaxation of internal stresses due to the differences in the properties, including volume mismatch, of oxide films and the substrate.
3.2. Dislocation arrangement along the oxide–substrate interface Further details of TEM observations in a plane parallel to the scale–matrix interface are shown in Fig. 3. Although, in all the samples, whether annealed in vacuum or in air, the deformation is macroscopically uniform, areas with highly nonuniform dislocation density are observed. This implies that local stress concentrators are present in the scale or in the region near the scale–matrix interface. The dislocation density changes drastically from a very low level, of 1012 m − 2, up to an upper limit of 1015 –1016 m − 2 along a short distance of several microns (see Fig. 3). At this stage of the investigations no quantitative analysis of the density of these areas was attempted. It has however been found that their density in the thin foils prepared from air-annealed samples is significantly higher than that observed in vacuum-annealed samples.
3.3. Microstructure of the oxide scale In order to explain how the scale influences the plastic deformation of specimens an attempt was made to characterise the microstructures of the oxides developed as a result of annealing at 600°C in air and vacuum. Specimens for TEM observations of the crosssection of the scale developed in vacuum are relatively easy to prepare, whereas the specimens of the scale
Fig. 3. An example of the nonuniform dislocation density in a plane parallel to the scale – matrix interface in the samples of 316 stainless steel annealed in vacuum and in air.
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
95
At this stage of investigations one may conclude that the tough and perfectly dense continuous outer layer of the scale formed on the samples annealed in vacuum improves their mechanical properties compared to the scale on the samples annealed in air, which exhibit a poor adhesion to the metal substrate.
4. Conclusions
Fig. 4. A cross-section TEM image of the scale on 316 stainless steel annealed in vacuum.
produced in air appeared to be extremely difficult because of its poor adhesion to the matrix. Fig. 4 shows a montage of TEM images of the cross-section of a vacuum-deposited scale. In the upper part of the figure an amorphous epoxy layer (used for gluing) is visible. Below the epoxy, there is a layer of uniform material with a submicron grain size of about 0.2 mm. The thickness of the layer is of the order of 1 mm. Beneath this layer, there is a thinner layer with an even smaller grain size. While the interface between these two layers is rather flat, the interface between the second layer and the metal substrate is profoundly rough. It is interesting to note that the layer next to the substrate seems to be porous, contrary to the outer layer, which is perfectly dense. The difficulty in preparing TEM specimens of the scale on the air-annealed samples results from its poor adhesion to the substrate. Fig. 5 shows the substrate (lower half of the picture) and the layer of scale with rather complicated structure (upper part of the picture). A part of the scale, where polycrystalline material is present, is similar in character to the upper layer of the scale formed on the samples annealed in vacuum.
The results obtained in the present work lead to the following conclusions. (1) The distribution of the dislocation density determined in the samples annealed in vacuum and in air suggests that the surface features affect the local plastic deformation and, thus, the mechanical properties of austenitic stainless steel. (2) The high dislocation density near the scale–matrix interface of strained specimens indicates the presence of the active dislocation sources in the vicinity of the free surface. (3) The oxide scale developed on the specimens after annealing at 600°C in vacuum consists of a layer of perfectly dense, submicron polycrystalline material. (4) The oxide scale developed on the specimens after annealing at 600°C in air is characterised by poor adhesion to the substrate. (5) The higher value of the flow stress determined in the vacuum-annealed specimens may perhaps be attributed to the high strength and good adhesion of the oxide scale.
Acknowledgements This work was supported by the State Committee for Scientific Research (KBN) Grant No. 7 T08A 008 12.
Fig. 5. A TEM image of a cross-section of the scale formed on 316 stainless steel annealed in air.
96
W. Zielin´ski et al. / Materials Science and Engineering A249 (1998) 91–96
References [1] R.M. Latanision, J.T. Fourie (Eds.), Surface Effects in Crystal Plasticity, Noordhoff, Leyden, 1977. [2] Viewpoint set No. 18 on: Surface Effects in Cyclic Deformation and Fatigue, Scripta Metall. Mater. 26 (1992) 1499. [3] K.G. Samuel, S.L. Mannan, Rodriguez, Acta Metall. 8 (1988) 2323. [4] A.W. Sleeswyk, Acta Metall. 6 (1958) 599. [5] C.F. Jenkins, G.V. Smith, Trans. Metall. Soc. AIME 245 (1969) 2149. [6] I.S. Kim, M.C. Chaturvedi, Metals Sci. 13 (1979) 691. [7] E.O. Hall, Yield Point Phenomena in Metals and Alloys, Macmillan (Eds.), UK, 1970.
.
[8] T. Nambu, J. Fukumori, M. Morinaga, Y. Matsumoto, T. Sakaki, Scripta Metall. Mater. 32 (1995) 407. [9] M. Morinaga, Y. Murata, M. Furui, T. Wada, Scripta Metall. Mater. 6 (1997) 699. [10] A.A. Abduluyahed, K. Rozniatowski, K.J. Kurzydłowski, Scripta Metall. Mater. 33 (1995) 1489. [11] A.A. Abduluyahed, K. Sikorski, K. Rozniatowski, K.J. Kurzydłowski, in: F.W. Poulsen, N. Bonanos, S. Linderoth, M. Mogensen, B. Zachav-Christiansen (Eds.), Proceedings of 17th Riso Int. Symp. on Mater. Sci., Roskilde, Denmark, 1996, pp. 155 – 160. [12] K.J. Kurzydłowski, Mater. Sci. Eng. A234 – 236 (1997) 1083. [13] K.J. Kurzydłowski, B. Ralph, Quantitative Description of Microstructure of Materials, CRC Press, Boca Raton, FL, 1995.