102

may cause the carbides to eoneentrate in the intergranular boundary region.

2. The presence in the steel of carbide-forming ele- ments slows down the redistribution of the carbide phase. The presence of non-carbide forming elements (e. g. , silicon), leads to intensive redistribution of the carbides in the direction of the iutracrystall ine boundaries.

3. Redistribution of the carbon in Cr-Mo-V-Si steel of the pearli te class is most intensive at 350 and 500 ~ Raising the temperature to 650 ~ leads to an appreciable weakening of the process .

REFERENCES

1. Bruk, B. I. Study of distribution of elements in welds by radiographic method. Welding, Coll. Articles No. 1, Sudprongiz, 1958.

2. Zemzin, V. N. "Boiler turbine construction", No. 6, 1950.

3. Bruk, B. I. Radioactive isotopes in metallurgy and metallography of welding, Sudpromgiz, 1959.

4. Yur'yev, S. F. and others. DAN SSSR, Vol. 4, No. 4, 1955.

5. Zav'yalov, A. S. and other. Radiographic study of carbon distribution in iron alloys, Metallography, Coil. Articles, Sudpromgiz, 1957.

6. Zav'yalov, A. S. and others . Laws governing intra- crystalline distribution of elements in metal alloys, Metallography, Coil. Articles No. 2, 1958; Physics of Metals and Metallography, Vol. 8, i ssue 3, 1959.

7. Bal'shin, M. Yu. Powder metallography, Metallur- gizdat, 1948.

8. Bokshteyn, S. Z. and others. Metallography and treatment of metals, No. 2, 1957.

9. Boksbteyn, S. Z. and others. Plant Laboratory, No. 3, Vol. 23, 1957.

10. Gertsriken, S. D. and others. Study of diffusion of cobalt and iron along grain boundaries, Study of cor ros iou- res i s tan t alloys, Vol. 4, USSR Academy of Sciences Press , 1959.

11. Zav'yalov, A. S. and others. P rocesses during tempering of alloyed steels, Metallography, Coll. Article No. 2, 1958.

12. Bruk, B. I. Solubility of carbon in alpha iron, Metallo- graphy, Coll. Articles No. 3, Sudpromgiz, 1959.

13. Bru]% B. I. DAN SSSR, Vol. 128, No. 4, 1959. 14. Petrova, Ye. F. andothers . Ibid.,Vol. 121, No. 6, 1958.

CHANGE IN DISLOCATION OF 1Kh18N9T STEEL

DURING HARDENING A N D SOFTENING

Eng. M. P. USIKOV AND L. M. UTEVSKIY, Technical Sciences Candidate

The authors studied the change in the dislocation s t ructure during the hardening and softening of 1Kh18N9T steel.

Rolled specimens 0.04 mm thick were stretched 0.3-- 10% after being annealed in evacuated ampoules at 1100 ~ (failure usually occurred after an elongation of 8 or 10%).

Some of the un-annealed specimens with a 96% defor- mation were heated at 400--800 ~ for an hour.

After the final heat t reatment and deformation, the s tr ip was thinned down by electrolytic polishing in a mixture of 60% H3PO 4 and 40% H2SO4 at a current density of 2--4 a /cm2 and at 60 ~ [1] .

The resulting foil 1000--2000 • thick was inspected with a UEMB--100 electron microscope at an accelerating voltage of 75 kv; the ~tiameter of the illuminating beam was not more than 10 microns .

The dislocations in the t ransmiss ion electron micro- scope produce an image on account of the fact that the distorted zone (usually on one side) scat ters the electrons to a greater extent along the dislocations, and the intensity of the beam pass ing through this zone is reduced. Hence, in a bright-field image the dislocations show up in the form of dark lines.

It has been shown both experimentally and theoretically [2] , that the width of the dislocation image (in either

s c r e w o r edge dislocation) ranges f rom 100 to 200 A. This makes it possible to distinguish separate dislocations where their density is about 1011 -- 1012 cm-2.

The visible dislocation line almost coincides with the projection of the actual dislocation onto the plane of the image. One end of the dislocation usually lies on the top surface and the other end on the bottom surface of the foil. The arrangement of the dislocation in the foil is shown in Fig. 1. Clearly, the length of the projection increases as the angle between the surface of the foil and the slip plane in which the dislocation lies is reduced.

Fig. 1. Arrangement of dislocation (1 and 2) on slip plane (3) in metal foil

Fig. 2 shows the s t ructure of an annealed specimen. For practical purposes there are no dislocations inside the grains. On the sloping twin boundary (Fig. 2) can be seen a hexagonal dislocation network showing slight deviation in disorientation f rom strictly twin disorientation. The row of dark bands along the boundary are due to diffraction and are lines of equal thickness on a wedge-shaped boundary. Slight deformation leads to the appearance within the grain of separate dislocations unevenly distributed on account

103

Fig. 2. Stainless steel. Annealing at 1100% One coherent twin boundary contains hexagonal dislocation network. • 47,000

of the different orientation of the grains with respect to the axis of tension. The density of the dislocations is slight and there is hardly any interaction between them (Fig. 3).

deformation as well); the mater ia l then begins to ha~cden. The dislocation density for 2% deformation is 2 �9 l0 w cm-2.

As the degree of deformation is increased, the dis- location s t ructure becomes more complex and there begin to appear intersections of dislocations, i r regu la r networks due to interaction between dislocations in different slip sys tems (Figs. 4a and b). Non-uniformity in the distribution of the dislocations is reduced.

After 96% deformation, the dislocation s t ructure is so complicated that it is not possible to pick out the individual elements (Fig. 4c). Thick clusters of dislocations fill the entire volume of the specimen. The grain and block bound- ar ies cannot be sharply resolved. The s t ructure is evi- dently fur ther complicated by a slight amount of martensi te (5--10%) forming during plastic deformation1.

Thus, in the annealed material the dislocation density is so smal l that it cannot be detected at all inside the grain; the field of vision of a microscope with 15, 000 magnification is 25 microns, while at a dislocation density of approxi- mately 105 cm -2, there is 1 dislocation per 1000 microns2. The dislocations are only observed on the grain boundaries (of blocks or twins) with slight disorientation.

Finally, considerable deformation (more than 8%) is accompanied by the formation of numerous dislocation networks; thick "clouds" of dislocations occur within the grains and make plastic deformation very difficult. The dislocation density is 1010 cm -2 for 10% deformation. Highly characterist ic is the state after cold rolling (96%); here the dislocation density is difficult to determine, but appears to be more than 1012 cm-2.

Heating to 600 ~ does not resul t in any appreciable r e - finement of the s tructure. Softening when the specimens are heated until recrystallization begins does not lead to any visible change in the arrangement or density of the dislocations. At 625 ~ however, recrystall ization begins (Fig. 5a), sub-grains appear, f ree f rom dislocations and with clear-cut boundaries. A further r i se in temperature

Slight plastic deformation (up to 1%) leads to the appear- ance of separate dislocations within the grain; the dislocation density calculated for several microphotographs attains about 109 cm -2.

There is only one slip system in action in each grain, and the dislocations are able to move about the grain un- hampered, except at the boundaries. This state corresponds to "easy glide" before hardening has occurred.

During deformation of more than 1%, there are secondary sys tems of slip. We observe interaction between dislocations moving in different slip planes, as a resul t of which there may form networks, and Lomer-Cottrel l b a r r i e r s obstructing the motion of the dislocation (and consequently fur ther plastic

1 Fig. 3. Stainless steel. 1% deformation, individual After 10% tensile deformation no martensi te was observed, dislocations within the grain, x 32, 000

104

produces growth of the subgrains and an increase in the dislocation-free volume of the specimen (Fig. 5b). At 700 ~ p r imary recrystall ization is complete and there are no dislocations within the grains. In this state the s t ructure only differs f rom the annealed state (at high temperatures) by having a smal le r grain size (Fig. 5e).

Thus, our research confirmed that the hardening and softening p rocesses in metal are accompanied by successive

changes in the dislocation structure.

REFERENCES

1. Kelly, P. M., Nutting J . , "Journal Institut of Metals", VIII, 1959.

2. Whelan, M. J . , ibid.

Fig. 4. Variation in s t ructure of stainless steel during in- crease in degree of deformation: a) 2% deformation, x 3400; b) 10% deformation, x 32, 000; c) 96% deformation, x40, 000; There are i r regu la r networks of dislocations inside the grains

Fig. 5. Variation in s t ructure of stainless steel when heated after deformation: a) at 625 ~ x 27,000; b) at 650 ~ • 30, 000;

c) at 700 ~ . x30,000