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A new tensile test on in situ solidified notched specimens: hot ductility analysis of continuous casting steels
Journal of Materials Processing Technology, 32 (1992) 325-334
325
Elsevier
A new tensile test on in situ solidified notched specimens: hot ductility analysis of continuous casting steels P. Deprez, J.P. Bricout and J. Oudin Mechanical Engineering Laboratory, C.N.R.S. Research Unit D 1401, MECAMAT, Universit6 de Valenciennes et du Hainaut Cambr6sis, BP 311, 59304 Valenciennes Cedex, France Abstract In order to securely predict the hot ductility and prevent hot-cracking propensity of continuous casting steels, a new tensile test device available for in situ solidified notched specimens is proposed. The test analysis makes it p o s s i b l e to d e f i n e the t h e r m o m e c h a n i c a l properties d i r e c t l y after solidification, without the usual cooling step down to room temperature and the following one to the test temperature. One of the original feature of the test consists in the shifting of the shrinkage cavity toward the unstrained region of the specimen and this allows accurate measurements in the notch non-segregated zone as well as secure observations of fracture types.Tests of two reference steels differing by their solidification mode show that the temperature path is a very significant parameter of the hot ductility. 1. I N T R O D U C T I O N The thermomechanical properties of steels are still poorly known at hot temperatures immediately after solidification. So it becomes important for the improvement of the continuous casters, to state precisely such properties and mainly the hot ductility after solidification in order to avoid hot defects occurrences such hot tears and brittle cracks during further straightening and hot w o r k i n g . The final m e t a l l u r g i c a l structure m a i n l y depends on temperature path and so far the hot ductility also. Up to now, the influence of the temperature path has received relatively little attention. One of the related problems is to find a way of simulating a solidification procedure and at the same time perform tensile tests at the required temperatures. So, the hot ductility of steels is usually defined from conventional tensile tests 0924-0136/92/$05.00 © 1992 Elsevier Science Publishers B.V. All rights reserved.
326 performed on tensile specimens in the conditions described in Figure la. It is supposed that the holding at 1350°C leads to a structure closely similar to the as-cast condition : coarse grains, dissolution of precipitates. In fact the microstructure at the testing temperature obtained in this way may differ from the microstructure obtained if the steel is first allowed to solidify and then to cool to testing temperature directly [1,2]. The differences in microstructure can entail that the steel will display different hot ductility performance. The present investigation was undertaken to develop a new tensile test on in situ solidified notched specimens able to reproduce the metallurgical structure after solidification according to a given temperature path as for instance the one shown on Figure lb. Using this tensile test, the influence of the temperature path on the hot ductility of two reference low carbon steels which differ by their solidification mode was considered for temperatures ranging from 750 to 1300 °C.
T (°C)
(a)
fusion
(b)
fusion
(c)
TIME
Figure 1. Temperature paths : (a) standard tensile test, (b) F-test, new tensile test on in situ solidified specimens, (c) R-test, tensile test on heated specimens from room temperature.
2. T E N S I L E T E S T D E V I C E The starting specimen-set consists in two round-bar blanks and a notched silica crucible; the specimen set and crucible are placed inside an inductor and then the two blanks are m o l t e n in a flowing argon a t m o s p h e r e (Figure 2a). The progressive rising of the testing machine lower ram allows the c r u c i b l e to fill up (Figure 2b); the crucible is d e s t r o y e d after solidification of the specimen owing to a clip device then the tensile test is achieved at the required temperature which is real time-controlled by a germanium photodiode pyrometer (Figure 2c).The choice of the specimen
327 dimensions has been achieved so as to shift the shrinkage cavity out of the notched zone (Figure 3). It is a basic advantage for accurate measurements in the non-segregated zones as well as secure observations of fracture types. Moreover, the tensile test of notched specimens offers a great interest due to negative values of hydrostatic pressure occurrences favouring the decohesion and related defects [3].
: fixedra>'-~;/'//m"~ 7////////////~ upperhalf ~.J l blank - ~ hlo • inductor--~ ~ ) ~ l ~ • • silica ____._.~(rL-~)) • • crucible III III lowerh a ~ ' - ¢ - - ' - ~ blank - "~///..///////,~ (a) initial specimen
shrinkag~
;v ,-T_LI o •
• (b) fusion
(o
O IZ--~I O (c) tensile test
Figure 2. Test procedure for in situ solidified notched specimens.
R=2
molten zone
Figure 3. In situ solidified specimen, main dimensions. 3. T E M P E R A T U R E P A T H I N F L U E N C E The t h e r m o m e c h a n i c a l features of steels are d e p e n d i n g on the solidification structure, the chemical composition heterogeneity and on the insoluble impurities distribution. So, in these conditions, it is obvious that reheating after solidification and cooling may have an influence on the intrinsic hot ductility of steel by modifying the structure and the location of the i m p u r i t i e s . The m e t a l l u r g i c a l p h e n o m e n a such as the y ---> o~ t r a n s f o r m a t i o n , the d y n a m i c recrystallisation of austenite and the
328
precipitation occurring in the poor hot ductility region at temperatures ranging from about 750 to 1100 °C are also affected by the temperature path. In order to illustrate the temperature path influence on hot ductility, tests were performed on two reference low carbon steels (Table 1) which differ by their solidification mode. Table 1 Chemical composition of A-steel and of B-steel (w%) C
Mn
P
S
Si
AI
Ni
Cr
Nb
V
N
Asteel
0.174 0.625 0.011
0.020 0.215 0.000 0.129 0.167 0.000 0.001 0.007
Bsteel
0.051 1.157 0.008 0.0004 0.294 0.026 0.020 0.015 0.042 0.048 0.006
The round-bar blanks were heated at 10 °C.s-1 up to 1520 °C for melting. After a lmin. holding, the specimen was solidified and cooled at 100°C.min -~ then strained until rupture, the elongation rate being 25 mm.min -~ and the tensile test temperatures ranging from 750 to 1300 °C. For each test, values of the peak load-initial notch cross section ratio, total elongation and notch diameter were recorded as strength and hot ductility parameters. Figure 4 shows the A-steel and B-steel typical hot ductility curves in terms of reduction RA of notch cross section versus test temperature; in so-called Ftests, the specimen is tested immediately after solidification (Figure lb) and in R-tests, the specimen is in situ solidified first, cooled at room temperature second, heated at 1350 °C during 1.5 min third and fourth tested at the required temperature (Figure lc). It can be noticed that it is difficult to foresee a general behaviour about temperature path influence on hot ductility of A-steel and of B-steel. At temperatures ranging from about 900 to 1100 °C, the F-tests with A-steel give a lower hot ductility than the one deduced from R-tests ; on the other hand, the B-steel shows an opposite behaviour. The gaps between the F/R-tests are both rather significant which gives evidence of the temperature path consequence as for the hot ductility of steels. The effect of the carbon content on hot ductility is particularly significant in the case of steels directly tested from solidification [4] : the intergranular cracking susceptibility is maximum for carbon contents about 0.18%. This effect seems correlated to a net y grain growth during the processes of solidification and subsequent cooling. For A-steel (0.17%C), the interdendritic segregations occurring during solidification and cooling are essentially localised on y grains boundaries (Figure 5) and lead to a significant embrittlement ; on the other hand these segregations are partly redistributed inside the new austenitic grains if reheating is achieved, so steel becomes less sensitive to intergranular cracking.
329
100-
RA%
80" 60" 40.
F-tests
40-
_~,~4
/
~
F-tests
20
20. A steel
0
!
800
(a)
w
|
1000
'
12'n0v
B steel , 0 700 , 8 ; 0 , i , 10~00 ' ~ ' 12~00 , , (b) T E S T T E M P E R A T U R E , °C
Figure 4. A-steel and B-steel hot ductility curves from F-tests and R-tests at 25 mm.min -1 elongation rate.
The variations of peak load-initial notch cross section ratio values (Fmax/S0) versus temperature for A-steel are presented in Figure 6. It can be seen that the F-tests curve is located below the R-tests curve for temperatures below 1150 °C, indicating that the difference in hot ductility was due to differences in structural hardening especially caused by differences in grain size.The microphotographs related to the unstrained region of the tensile specimen have shown that the austenite grain of steels directly tested from solidification was indeed perceptibly larger than in R-tests case. These facts justify the widening to higher temperatures of the ductility trough for Ftests. Note that A-steel has a relatively high sulphur content and a Mn/S ratio of about 30, so an intergranular embrittlement due to a liquid grain boundary film, so called hot shortness, is possible; the FeS segregated at the austenite grain boundaries may be indeed liquid at temperatures above 980 °C. The existence of both a high temperature (1100 °C) for the beginning of the ductility trough and a discontinuity on F-tests hot ductility curve about 980°C confirms this assumption. This behaviour is not noticeable on the R-tests curve, so it seems that the temperature path fixed for R-tests (Figure lc) is less sensitive for the detection of this phenomenon that is considered, in addition, as a boundary case for the values of the Mn/S ratio of about 30 (Figure 7).
330
~.1540~'"m -31520"1 -_.
_
liquid
I,x\\\x.x\x.\,~intergranular segregations
I
~ --~--...
, :<, ,oo~~,. ~1
.
~'tF_-¢, '~l~lf~
.
'14401 4'°1/1,'E,,'!, ~+)¢i-" ~-t ~, ! I I },i,'-
~,!
1420-11~ |
I'.~1
1
4
0
....
~l---- 0% Mn
I
Solidification
I
into Y phase into 8 phase (A-steel) (B-steel) SOLIDIFICATION MODES
~ i - o.~,%MnI 0
"
t
~
4
Solidification
CARBONE CONTENT, C%
Figure 5. Equilibrium Fe-C phase diagram and solidification modes.
100 80 ]
F-tests R-tests
¢f
LLs
60 4O 20
0
'
"-
' 1300 TEST TEMPERATURE, °C
Figure 6. A-steel, peak load-initial notch cross section ratio versus temperature for F-tests and for R-tests, elongation rate 25 mm.min -1. In the case of the niobium alloyed B-steel, the decay in RA is more progressive for the tests directly achieved after fusion and the minimum
331
value of RA is shifted about 100 °C toward low temperatures. Due to the low carbon content (0.051%), the solidification thoroughly takes place in the ~i domain without peritectic reaction occurrence, thereby 7 grains formation is independent on the solidification structure [5] ; segregations are thus located inside 7 grains (Figure 5).
" 1.6 14 "-"
/
" 1.2-
(Mn/S = 100) /
1.0-
~
ductility
-
7
A-steel
O.
oo. 0 0
0.01
0.02
Sulfur (wt %) Figure 7. Influence of Mn/S ratio on hot ductility of steels from [6], A-steel Mn-S contents location on the graph. Logically the F-tests and R-tests must give similar results however the ferrite formation kinetics and the ferrite morphology are probably different for the two respective temperature paths and it is likely that ferrite formation would be delayed in the F-test case including a continuous cooling from solidification. Moreover, it is remarkable to state that the increase of RA at low temperatures, usually associated with ferrite formation, is only noticed for R-tests about 800 °C. The embrittlement due to the niobium effect originates in the dynamic precipitation of carbonitrides such as Nb(CN) inducing a strengthening of austenite and a lower "ygrain boundary cohesion, moreover, the pinning effect of the Nb(CN) particles on ~, grain boundary migration makes recrystallisation rather difficult [7]. Figure 8 shows the effect of test temperature on peak load-initial notch cross section ratio values (Fmax/S0). It can be noticed that the F-tests curve is situated below the R-tests curve for temperatures less than 1000 °C. It is likely that the dynamic precipitation of carbonitrides Nb(CN) occurring during F-tests gives coarsened particles resulting in softening of the austenitic matrix. So the
332
dynamic recrystallisation of austenite might be easier for tests immediately performed after solidification. The microphotographs shown in Figure 9 confirm this assumption : the ferritic grain size in the vicinity of the fracture surface, stemmed from recrystallised austenitic grain, is indeed smaller for the F-test case. 200. 12. v
R-tests F-tests
* o
3
100
1
700
800
i
I
I
i
I
1000 1200 TEST TEMPERATURE, °C
Figure 8. B-steel, peak load-initial notch cross section ratio versus temperature for F-tests and R-tests, elongation rate 25 mm.min-1. The fracture surfaces not including the shrinkage cavity, it becomes possible to perform an objective SEM fractographic analysis. Figures 10 and 11 respectively show the effect of temperature path on the fracture surface of A-steel specimens tested at 1000 °C and B-steel specimens tested at 850 °C. A-steel exhibits some degree of intergranular fracture when it is tested after direct cooling from the fusion temperature and lots of intergranular precipitates are easily detected in the SEM fractograph (Figure 10a) whereas a ductile transgranular fracture mode is rather observed in the specimen tested after reheating (Figure 10b). On the contrary, the intergranular decohesion trend is prominent for the R-tests performed on B-steel specimens (Figure 11). 4. CONCLUSION A new tensile test on in situ solidified notched specimen has been achieved for investigations of steels thermomechanical properties immediately after solidification. The shrinkage cavity is shifted toward the unstrained region of the specimen and this allows analysis in the notch non-segregated zone as well
333 as secure observations of fracture types. Comparing tests performed on reheated specimens and on in situ solidified specimens of two reference low carbon steels, significant changes of the hot ductility curves are enhanced. It is suggested that the differences depend on the distribution of the segregations, the ferrite formation kinetics and the dynamic recrystallisation occurrence of austenite.
Figure 9. B-steel, tensile test at 950 °C, microphotographs of the strained region: (a) F-test, (b) R-test.
Figure 10. A-steel, tensile test at 1000 °C, microphotographs of fracture surfaces: (a) F-test, (b) R-test.
334
Figure 11. B-steel, tensile test at 850 °C, microphotographs of fracture surfaces: (a) F-test, (b) R-test.
5. REFERENCES 1 G. A. Wilber, R. Batra, W. F. Savage and W. J. Childs, Metal. Trans., 6A (1975) 1727. 2 B. Rogberg, Scand. J. Metallurgy, 12 (1983) 51. 3 J. Oudin, Y. Ravalard, J.C. Gelin, G. Lacombe and T. Labarthe-Vacquier, Mat6riaux et Techniques, 11-12 (1988) 39. 4 Y. Maehara, K. Yasumoto and Y. Sugitani, Transactions I.S.I.J., 25 (1985) 1045. 5 G. Bernard, J.P. Birat and J.C. Humbert, Rev. de M6t., 7 (1978) 467. 6 H. Suzuki, T. Nishimura and S. Yamaguchi, Tetsu to Hagane, 65 (1979) 2038. 7 Y. Maehara, K. Nakai, K Yasumoto and T. Mishima, Transactions I.S.I.J., 28 (1988) 1021.
Acknowledgments The authors are grateful to C.N.R.S., R6gion Nord Pas de Calais, D616gation R6gionale du Minist~re de la Recherche, Minist~re de l'Education Nationale and Sollac for support in the above developments.
325
Elsevier
A new tensile test on in situ solidified notched specimens: hot ductility analysis of continuous casting steels P. Deprez, J.P. Bricout and J. Oudin Mechanical Engineering Laboratory, C.N.R.S. Research Unit D 1401, MECAMAT, Universit6 de Valenciennes et du Hainaut Cambr6sis, BP 311, 59304 Valenciennes Cedex, France Abstract In order to securely predict the hot ductility and prevent hot-cracking propensity of continuous casting steels, a new tensile test device available for in situ solidified notched specimens is proposed. The test analysis makes it p o s s i b l e to d e f i n e the t h e r m o m e c h a n i c a l properties d i r e c t l y after solidification, without the usual cooling step down to room temperature and the following one to the test temperature. One of the original feature of the test consists in the shifting of the shrinkage cavity toward the unstrained region of the specimen and this allows accurate measurements in the notch non-segregated zone as well as secure observations of fracture types.Tests of two reference steels differing by their solidification mode show that the temperature path is a very significant parameter of the hot ductility. 1. I N T R O D U C T I O N The thermomechanical properties of steels are still poorly known at hot temperatures immediately after solidification. So it becomes important for the improvement of the continuous casters, to state precisely such properties and mainly the hot ductility after solidification in order to avoid hot defects occurrences such hot tears and brittle cracks during further straightening and hot w o r k i n g . The final m e t a l l u r g i c a l structure m a i n l y depends on temperature path and so far the hot ductility also. Up to now, the influence of the temperature path has received relatively little attention. One of the related problems is to find a way of simulating a solidification procedure and at the same time perform tensile tests at the required temperatures. So, the hot ductility of steels is usually defined from conventional tensile tests 0924-0136/92/$05.00 © 1992 Elsevier Science Publishers B.V. All rights reserved.
326 performed on tensile specimens in the conditions described in Figure la. It is supposed that the holding at 1350°C leads to a structure closely similar to the as-cast condition : coarse grains, dissolution of precipitates. In fact the microstructure at the testing temperature obtained in this way may differ from the microstructure obtained if the steel is first allowed to solidify and then to cool to testing temperature directly [1,2]. The differences in microstructure can entail that the steel will display different hot ductility performance. The present investigation was undertaken to develop a new tensile test on in situ solidified notched specimens able to reproduce the metallurgical structure after solidification according to a given temperature path as for instance the one shown on Figure lb. Using this tensile test, the influence of the temperature path on the hot ductility of two reference low carbon steels which differ by their solidification mode was considered for temperatures ranging from 750 to 1300 °C.
T (°C)
(a)
fusion
(b)
fusion
(c)
TIME
Figure 1. Temperature paths : (a) standard tensile test, (b) F-test, new tensile test on in situ solidified specimens, (c) R-test, tensile test on heated specimens from room temperature.
2. T E N S I L E T E S T D E V I C E The starting specimen-set consists in two round-bar blanks and a notched silica crucible; the specimen set and crucible are placed inside an inductor and then the two blanks are m o l t e n in a flowing argon a t m o s p h e r e (Figure 2a). The progressive rising of the testing machine lower ram allows the c r u c i b l e to fill up (Figure 2b); the crucible is d e s t r o y e d after solidification of the specimen owing to a clip device then the tensile test is achieved at the required temperature which is real time-controlled by a germanium photodiode pyrometer (Figure 2c).The choice of the specimen
327 dimensions has been achieved so as to shift the shrinkage cavity out of the notched zone (Figure 3). It is a basic advantage for accurate measurements in the non-segregated zones as well as secure observations of fracture types. Moreover, the tensile test of notched specimens offers a great interest due to negative values of hydrostatic pressure occurrences favouring the decohesion and related defects [3].
: fixedra>'-~;/'//m"~ 7////////////~ upperhalf ~.J l blank - ~ hlo • inductor--~ ~ ) ~ l ~ • • silica ____._.~(rL-~)) • • crucible III III lowerh a ~ ' - ¢ - - ' - ~ blank - "~///..///////,~ (a) initial specimen
shrinkag~
;v ,-T_LI o •
• (b) fusion
(o
O IZ--~I O (c) tensile test
Figure 2. Test procedure for in situ solidified notched specimens.
R=2
molten zone
Figure 3. In situ solidified specimen, main dimensions. 3. T E M P E R A T U R E P A T H I N F L U E N C E The t h e r m o m e c h a n i c a l features of steels are d e p e n d i n g on the solidification structure, the chemical composition heterogeneity and on the insoluble impurities distribution. So, in these conditions, it is obvious that reheating after solidification and cooling may have an influence on the intrinsic hot ductility of steel by modifying the structure and the location of the i m p u r i t i e s . The m e t a l l u r g i c a l p h e n o m e n a such as the y ---> o~ t r a n s f o r m a t i o n , the d y n a m i c recrystallisation of austenite and the
328
precipitation occurring in the poor hot ductility region at temperatures ranging from about 750 to 1100 °C are also affected by the temperature path. In order to illustrate the temperature path influence on hot ductility, tests were performed on two reference low carbon steels (Table 1) which differ by their solidification mode. Table 1 Chemical composition of A-steel and of B-steel (w%) C
Mn
P
S
Si
AI
Ni
Cr
Nb
V
N
Asteel
0.174 0.625 0.011
0.020 0.215 0.000 0.129 0.167 0.000 0.001 0.007
Bsteel
0.051 1.157 0.008 0.0004 0.294 0.026 0.020 0.015 0.042 0.048 0.006
The round-bar blanks were heated at 10 °C.s-1 up to 1520 °C for melting. After a lmin. holding, the specimen was solidified and cooled at 100°C.min -~ then strained until rupture, the elongation rate being 25 mm.min -~ and the tensile test temperatures ranging from 750 to 1300 °C. For each test, values of the peak load-initial notch cross section ratio, total elongation and notch diameter were recorded as strength and hot ductility parameters. Figure 4 shows the A-steel and B-steel typical hot ductility curves in terms of reduction RA of notch cross section versus test temperature; in so-called Ftests, the specimen is tested immediately after solidification (Figure lb) and in R-tests, the specimen is in situ solidified first, cooled at room temperature second, heated at 1350 °C during 1.5 min third and fourth tested at the required temperature (Figure lc). It can be noticed that it is difficult to foresee a general behaviour about temperature path influence on hot ductility of A-steel and of B-steel. At temperatures ranging from about 900 to 1100 °C, the F-tests with A-steel give a lower hot ductility than the one deduced from R-tests ; on the other hand, the B-steel shows an opposite behaviour. The gaps between the F/R-tests are both rather significant which gives evidence of the temperature path consequence as for the hot ductility of steels. The effect of the carbon content on hot ductility is particularly significant in the case of steels directly tested from solidification [4] : the intergranular cracking susceptibility is maximum for carbon contents about 0.18%. This effect seems correlated to a net y grain growth during the processes of solidification and subsequent cooling. For A-steel (0.17%C), the interdendritic segregations occurring during solidification and cooling are essentially localised on y grains boundaries (Figure 5) and lead to a significant embrittlement ; on the other hand these segregations are partly redistributed inside the new austenitic grains if reheating is achieved, so steel becomes less sensitive to intergranular cracking.
329
100-
RA%
80" 60" 40.
F-tests
40-
_~,~4
/
~
F-tests
20
20. A steel
0
!
800
(a)
w
|
1000
'
12'n0v
B steel , 0 700 , 8 ; 0 , i , 10~00 ' ~ ' 12~00 , , (b) T E S T T E M P E R A T U R E , °C
Figure 4. A-steel and B-steel hot ductility curves from F-tests and R-tests at 25 mm.min -1 elongation rate.
The variations of peak load-initial notch cross section ratio values (Fmax/S0) versus temperature for A-steel are presented in Figure 6. It can be seen that the F-tests curve is located below the R-tests curve for temperatures below 1150 °C, indicating that the difference in hot ductility was due to differences in structural hardening especially caused by differences in grain size.The microphotographs related to the unstrained region of the tensile specimen have shown that the austenite grain of steels directly tested from solidification was indeed perceptibly larger than in R-tests case. These facts justify the widening to higher temperatures of the ductility trough for Ftests. Note that A-steel has a relatively high sulphur content and a Mn/S ratio of about 30, so an intergranular embrittlement due to a liquid grain boundary film, so called hot shortness, is possible; the FeS segregated at the austenite grain boundaries may be indeed liquid at temperatures above 980 °C. The existence of both a high temperature (1100 °C) for the beginning of the ductility trough and a discontinuity on F-tests hot ductility curve about 980°C confirms this assumption. This behaviour is not noticeable on the R-tests curve, so it seems that the temperature path fixed for R-tests (Figure lc) is less sensitive for the detection of this phenomenon that is considered, in addition, as a boundary case for the values of the Mn/S ratio of about 30 (Figure 7).
330
~.1540~'"m -31520"1 -_.
_
liquid
I,x\\\x.x\x.\,~intergranular segregations
I
~ --~--...
, :<, ,oo~~,. ~1
.
~'tF_-¢, '~l~lf~
.
'14401 4'°1/1,'E,,'!, ~+)¢i-" ~-t ~, ! I I },i,'-
~,!
1420-11~ |
I'.~1
1
4
0
....
~l---- 0% Mn
I
Solidification
I
into Y phase into 8 phase (A-steel) (B-steel) SOLIDIFICATION MODES
~ i - o.~,%MnI 0
"
t
~
4
Solidification
CARBONE CONTENT, C%
Figure 5. Equilibrium Fe-C phase diagram and solidification modes.
100 80 ]
F-tests R-tests
¢f
LLs
60 4O 20
0
'
"-
' 1300 TEST TEMPERATURE, °C
Figure 6. A-steel, peak load-initial notch cross section ratio versus temperature for F-tests and for R-tests, elongation rate 25 mm.min -1. In the case of the niobium alloyed B-steel, the decay in RA is more progressive for the tests directly achieved after fusion and the minimum
331
value of RA is shifted about 100 °C toward low temperatures. Due to the low carbon content (0.051%), the solidification thoroughly takes place in the ~i domain without peritectic reaction occurrence, thereby 7 grains formation is independent on the solidification structure [5] ; segregations are thus located inside 7 grains (Figure 5).
" 1.6 14 "-"
/
" 1.2-
(Mn/S = 100) /
1.0-
~
ductility
-
7
A-steel
O.
oo. 0 0
0.01
0.02
Sulfur (wt %) Figure 7. Influence of Mn/S ratio on hot ductility of steels from [6], A-steel Mn-S contents location on the graph. Logically the F-tests and R-tests must give similar results however the ferrite formation kinetics and the ferrite morphology are probably different for the two respective temperature paths and it is likely that ferrite formation would be delayed in the F-test case including a continuous cooling from solidification. Moreover, it is remarkable to state that the increase of RA at low temperatures, usually associated with ferrite formation, is only noticed for R-tests about 800 °C. The embrittlement due to the niobium effect originates in the dynamic precipitation of carbonitrides such as Nb(CN) inducing a strengthening of austenite and a lower "ygrain boundary cohesion, moreover, the pinning effect of the Nb(CN) particles on ~, grain boundary migration makes recrystallisation rather difficult [7]. Figure 8 shows the effect of test temperature on peak load-initial notch cross section ratio values (Fmax/S0). It can be noticed that the F-tests curve is situated below the R-tests curve for temperatures less than 1000 °C. It is likely that the dynamic precipitation of carbonitrides Nb(CN) occurring during F-tests gives coarsened particles resulting in softening of the austenitic matrix. So the
332
dynamic recrystallisation of austenite might be easier for tests immediately performed after solidification. The microphotographs shown in Figure 9 confirm this assumption : the ferritic grain size in the vicinity of the fracture surface, stemmed from recrystallised austenitic grain, is indeed smaller for the F-test case. 200. 12. v
R-tests F-tests
* o
3
100
1
700
800
i
I
I
i
I
1000 1200 TEST TEMPERATURE, °C
Figure 8. B-steel, peak load-initial notch cross section ratio versus temperature for F-tests and R-tests, elongation rate 25 mm.min-1. The fracture surfaces not including the shrinkage cavity, it becomes possible to perform an objective SEM fractographic analysis. Figures 10 and 11 respectively show the effect of temperature path on the fracture surface of A-steel specimens tested at 1000 °C and B-steel specimens tested at 850 °C. A-steel exhibits some degree of intergranular fracture when it is tested after direct cooling from the fusion temperature and lots of intergranular precipitates are easily detected in the SEM fractograph (Figure 10a) whereas a ductile transgranular fracture mode is rather observed in the specimen tested after reheating (Figure 10b). On the contrary, the intergranular decohesion trend is prominent for the R-tests performed on B-steel specimens (Figure 11). 4. CONCLUSION A new tensile test on in situ solidified notched specimen has been achieved for investigations of steels thermomechanical properties immediately after solidification. The shrinkage cavity is shifted toward the unstrained region of the specimen and this allows analysis in the notch non-segregated zone as well
333 as secure observations of fracture types. Comparing tests performed on reheated specimens and on in situ solidified specimens of two reference low carbon steels, significant changes of the hot ductility curves are enhanced. It is suggested that the differences depend on the distribution of the segregations, the ferrite formation kinetics and the dynamic recrystallisation occurrence of austenite.
Figure 9. B-steel, tensile test at 950 °C, microphotographs of the strained region: (a) F-test, (b) R-test.
Figure 10. A-steel, tensile test at 1000 °C, microphotographs of fracture surfaces: (a) F-test, (b) R-test.
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Figure 11. B-steel, tensile test at 850 °C, microphotographs of fracture surfaces: (a) F-test, (b) R-test.
5. REFERENCES 1 G. A. Wilber, R. Batra, W. F. Savage and W. J. Childs, Metal. Trans., 6A (1975) 1727. 2 B. Rogberg, Scand. J. Metallurgy, 12 (1983) 51. 3 J. Oudin, Y. Ravalard, J.C. Gelin, G. Lacombe and T. Labarthe-Vacquier, Mat6riaux et Techniques, 11-12 (1988) 39. 4 Y. Maehara, K. Yasumoto and Y. Sugitani, Transactions I.S.I.J., 25 (1985) 1045. 5 G. Bernard, J.P. Birat and J.C. Humbert, Rev. de M6t., 7 (1978) 467. 6 H. Suzuki, T. Nishimura and S. Yamaguchi, Tetsu to Hagane, 65 (1979) 2038. 7 Y. Maehara, K. Nakai, K Yasumoto and T. Mishima, Transactions I.S.I.J., 28 (1988) 1021.
Acknowledgments The authors are grateful to C.N.R.S., R6gion Nord Pas de Calais, D616gation R6gionale du Minist~re de la Recherche, Minist~re de l'Education Nationale and Sollac for support in the above developments.