$3, $4) were obtained by applying the heat treatments given in Table 2. In order to reveal the microstructure, the specimens were polished mechanically and ...
Jommnlot
Materials Processing Technology ELSEVIER
Journal of Materials Processing Technology 69 (1997) 238-246
Investigation of the casting microstructure of a Fe-base, C r - N i - M o duplex alloy M. Ceylan a,*, V. Kuzucu a, M. Aksoy b, i. Aksoy c, M. Kaplan d, M.M. Yildirim d a University of Firat, Faculty of Science and Arts, Department of Physics, Elazi~,, Turkey b University of Firat, Faculty of Engineering, Department of Mechanical Engineering, Ela=i~,, Turkey c University of ]ni~nii, Faculty of Science and Arts, Department of Physics, Malatya, Turke.v a University of Firat, Faculty of Technical Education, Department of Metallurgy, Ela:i~,, Turkey Received 13 January 1o96
Abstract
The microstructures of Fe-base cast duplex stainless steel specimens, which contain Cr, Ni, Me and other elements and have been subjected to various heat treatments, were examined with X-ray diffraction, thermal analysis and optical microscopy and their hardnesses were also measured. From the metallographic examinations, it was seen that these specimens included austenite, ferrite, some inter-metallic phases, grain boundary carbides and oxide inclusions. It has been determined that the amount of grain-boundary carbides decreased, and second-phases precipitated in ferrite, in specimens annealed at 950°C for 2 h and then cooled i.n water and liquid nitrogen. In addition it has been found by X-ray diffraction that the heat treatments applied did not change the types of the phases, but caused lattice deformations in the phases. Furthermore, beside ferrite and austenite, Laves, o', G, g, R, X, Fe3C, Fe3P, FeS2, MnS phases and M23C6, M7C3, M6C carbides were determined to be in the microstructure. DSC (differential scanning calorimeter) curves showed that some stoichiometry compounds of small amounts occurred between 370 and 470"C, and in specimens cooled in liquid nitrogen after annealed at 950"C for 2 h, a phase with a small percentage dissolving at about 310"C. DTA (differential thermal analysis) curves also showed that austenite phase transformed to ferrite phase completely between 1182 and 1220"C. In addition, it was found that the hardness of the specimens were different. © 1997 Elsevier Science S.A. Keywords: Casting; Microstructure; Fe-base; Duplex alloy
1. Introduction
Stainless steels, which are a large group of ferrous alloys, are used against corrosion, wear and high temperature effects in corrosive environments [1]. The common types of these steels, which contain various alloying elements such as Cr, Ni, Me, V, W, C, Ti etc. in different proportions, are martensitic stainless steels, ferritic stainless steels, austenitic stainless steels, precipitating hardened stainless steels and ferritic-austenitic duplex stainless steels [2]. The specification of the austenitic-ferritie duplex stainless steel investigated in this study is F e - 2 5 . 5 C r 4.8Ni- 1.40Mo-0.83Mn-0.46Si-0.03C-0.02P-0.02S. Although duplex stainless steels were developed in the * Corresponding author. Fax: + 424 2330062. 0924-0136/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S0924rO I 36(97)O0025.3
1940s, they were first commercially produced in the 1970s [3]. Since these steels have the good properties of both ferritic and austenitic stainless steels, they have found an important application area in the last decade. They have a corrosion resistance close to that of ferritic steels. Their toughness is better than that of ferritic steels but worse than that of austenitic steels. Their strengths are higher than those of austenitic steels. Therefore, it provides an optimum material selection to prefer these steels for industrial requirements [4]. When the ternary F e - C r - N i equilibrium diagram is studied, the importance of heat treatment and alloying elements are understood in the formation of 0c and ~, phases. When this material is annealed at 10001100°C, it has a homogen duplex structure and there is no precipitated phase. In general, rapid cooling is necessary to prevent the precipitation of other phases [4].
M. Ceylon et al./Journal of Materials Processing Technology 69 (1997) 238-246
239
Table ! Weight percentages of the alloying elements Fe
Cr
Ni
Mo
Mn
Si
C
S
P
68.35
25.03
3.88
0.89
0.82
0.77
0.19
0.027
0.027
On the other hand, because these steels are not stable under 1000°C, the formation of chromium-rich a, ;~, R etc. phases are expected. However, these inter-metallic phases are not desired. The inter-metallic a phase in austenitic-ferritic alloys is important, because the presence of a reduces the corrosion resistance and imparts ambient-temperature brittleness [5]. In this study, the microstructures of: (i) the specimen obtained from casting; (ii) the specimen obtained by cooling in water after being annealed at 950°C for 2 h, which is insufficient for homogenizing; (iii) and (iv) the specimens obtained by cooling the specimens above, respectively, in liquid nitrogen from room temperature were examined with X-ray diffraction, the optical metal microscope and the DSC and DTA methods. The hardness of these four specimens was measured on the Vickers scale.
with modified Murakami's etchant (30 g KOH + 30 g K 3 F e ( C N ) 6 + 150 g H 2 0 ) , whereupon the precipitated phases became coloured. Colour and black and white photographs of the microstructures of the specimens were taken with an Olympus metal microscope, in some photographs inversionat oil being used. Powder samples were taken from the specimens in Table 2 for X-ray diffraction and theral analysis. X-ray diffractograms for each specimen were taken using a Rigaku-Geigerflex diffractometer. DTA (differential thermal analysis) and DSC (differential scanning calorimeter) results were obtained from using a Shi-
2. Experimental Twenty five kg of the mixture was melted in an induction furnace, cast into sand moulds having the dimensions 70 x 40 x 300 mm and then left to cool. The contents of alloying elements were determined, as shown in Table 1, and four different specimens (S1, $2, $3, $4) were obtained by applying the heat treatments given in Table 2. In order to reveal the microstructure, the specimens were polished mechanically and then etched to obtain the black and white appearance of the microstructure with HNO3 + C2H402 + 4HCI + H20 solution. In addition, in order to distinguish the precipitated phases, they were etched for a second time at room temperature
A
B
C
D
Table 2 Heat treatments applied to the specimens and the symbols of specimens The symbol of sample
Heat treatment
Sl S2
The cast specimen SI cooled in water after being annealed at 950°C for 2 h SI cooled in liquid nitrogen for 7 min $2 cooled in liquid nitrogen for 7 min
$3 $3
Fig. 1. Surface photographs of specimens: (a) SI ( x 100); (b) $2 ( x 100), (c) $3 ( x 200); and (d) $4 ( x 200).
¸240¸~
Mi Ceylan et ali ~Journal of Materials Processing Technology 69(1997)238-246
Fig. 3. A colour surface photograph of specimen Sl ( x 2000). Fig. 2. The needle,like morphology of some of the austenite in specimen SI ( x 665).
madzu 50 instrument using the powders of the same samples. In addition, the hardness of the specimens was measured on the Vickers hardness scale using a load of 5 kg, as shOWn in Table 2.
3. Results and discussion
3.1. Evaluation o f the microstructure Surface photographs of specimens S 1- $4 etched with HNO3 + C 2 H 4 0 2 @ 4HC1 + H20 etchant are shown in
Fig. 4. Colour surface photographs of specimen $2: (a) and (b) show M7C3, M23C6 and M6C carbides, ~,-Fe, 0~-Fe, a phase, L (Laves) phase and ! (inclusion ( x 2000); (c) fine dispersion of spheridai carbides (M6C) in a matrix of ferrite ( x 2000); and (d) a phase (blue coloured) and other precipitates in u-ferrite ( × 1000).
M. Ceylan et al./Journal of Mater&Is Process&g Technology 69 (1997) 238-246
.
3
4
S
6
7
241
9
!0
$2
$3
_Ad (a)
(b) Fig. 5. Colour surface photographs of specimen $3: (a) ( × 2000) and (b) show ( x 401)) show M,~C6 and M7C3 carbides, 7-Fe, ~-Fe, a phase and inclusion (1).
Fig. l(a-d), respectively. In all of specimens, the matrix is ferrite and the second phase is austenite. The black phases in the grain boundaries are chromium-rich M23C6 metal carbides, where M represents one or more types of metal atom. As can be seen from the photographs, heat treatments did not change the amounts and the kinds of the phases, but in specimens $2 and $4, which were annealed at 950°C for 2 h, the grain
Fig. 6. Colour surface photograph of specimen $4 ( × 2000).
='
Fig. 7. X-ray diffractoms of specimens S1-$4.
boundary carbides did not disappear completely, despite the decrease in their amount, because of the insufficiency in the annealing temperature [3]. Some of the austenite grains in the ferrite matrix has an acicular or needle-like morphology, which can be seen clearly in Fig. 2. This kind of morphology can be achieved with a slow cooling from temperatures above the austenite solvus line, with a cooling rate similar to that of air cooling [6]. However, even in the case of annealing first and then cooling in liquid nitrogen, some amount of austenite retained this morphology, as the structure of the austenite at room temperature did not change at 950°C (F;.g. !). Etching with modified Murakami's etchant at room temperature followed by the etching with the etchant above enabled the phases to be clarified easier by giving the phases different colours. The modified Murakami's etchant used especially to reveal the a phase, makes the ferrite phase yellow and the a phase blue, austenite either being not affected or having a pale yellow colour [5]. This etchant can be used to reveal metal carbides and the other precipitated phases. In Figs. 3-6, surface photographs of specimens S1-$4 etched with modified Murakami's etchant are shown. Except for Fig. 4(d) and Fig. 5(b), all of the photographs were obtained using inversional oil, the ~r phase revealed in blue normally--being seen as pale brown because of the difference in the refraction index of the oil. The tr phase, which is hard and brittle, generally precipitated on carbide particles on grain boundaries in large blocks. It was decided that the red phases seen in the photographs were oxide inclusions, these inclusions not being affected by heat treatment. Black phases seen on the grain boundaries are chromium-rich M23C6 metal carbides, these carbides decreasing by a significant amount in annealed $2 and $3 specimens. On the other hand, in Fig. 3, M6C carbides are seen as distributed in the matrix in addition to M23C6 carbides and oxide inclusions. A part of these carbides precipitated on grain boundaries. In Fig. 4(a) and (b), Laves phase and MTC 3 carbides are also seen together with the a phase,
M. Ceylan et aL /Journal of Materials Processing Technology 69 (1997) 238-246
242
Table 3 The list of 20 angles and mean values of 20 angles used in calculations Peak No.
20 for Si
20 for $2
20 for $3
20 for $4
Mean values of 20
1 2 3 4 5 6 7 8 9 I0
18.14 20.47 23.04 25.57 28.88 31.43 31.67 35.62 41.15 46.31
18.27 20.55 23.16 25.73 28.99 31.63 -35.69 41.29 46.37
18.16 20.47 23.04 -29.01 31.47 -35.64 41.33 46.27
17.97 20.28 22.89 25.49 28.87 31.33 32.94 35.37 40.99 4.17
18.13 20.44 23.03 25.60 28.94 31.46 32.30 35.58 41.19 46.28
Table 4 The crystal lattice planes which contribute to the X-ray peaks of expected phases Phases
Peak No. I
• -Fe y-Fe Mo FeCr2S4
2
3
110 I11
200
4
044
004
402
005
411
510
440
404
443
543 550
M23C6
400 002 500 430 332
442
Z (Chi) phase
400
333 511 420
a-phase
002
Laves phase
111 Ell 113 2i3 123 224 134 002 030
G phase
It phase (FevMo6)
R phase Fe3C Fe7C3 FeS,
211
P M7C3 (Cr7C3)
422 420
M6C MnS
224
202 022 201 221 223 202
500 430
222
235
055 445
355 445 022 005 034
8
9
l0
211
220
310 400 222
540 203
013 125
442 555
544
554
521
532
542
402
123
442
433
404
014 114
/24
233
310 225 224
4744
226
550 543
306
323
332
112 031 002 300
122
212 032
221 300 443 322 122 431 424
311
440 002 421 430 404
7
220
005 034
Fe~P
6
200
110 024 133
fl-Cr • -Mn
5
410 5T0
524
444
535
544
556
602
113
203
400
332
333
005 324
322 410 555 003 442
303 411
432
442
660 400
555
232
444
553
355
123
224 504
420
775 333 511
777
M. Cevkm el al. / Journal of Materials Processing Teclmology 69 (1997) 238.- 246
243
Table 5 The crystal structures, lattice parameters and f o r m u l a s o f expected phases Phase
Crystal structure
Lattice parameter (A)
Formula
Reference
x-Fe 7-Fe Mo FeCr_,S4 //-Cr :e-Mn FeaP G-phase M_,sC~,
bcc fcc bcc
a = 2.8665 a = 3.6467 a = 3.1470
:~-Fe ;'-Fe
[10] [BO]
Mo
[ 10 l
[7.,('r S
fcc bcc tetragonal cubic fcc
a a a a a
= = = ~ =
3.86 8.9126 9.107, c = 4.460 11.3 10.57-10.68
varies with composition /. phase
bcc
a = 8.862-9.188
varies with composition a phase
tetragonal
Laves phase
hegzagonal
p phase R phase
hexagonal
FesC
orthorhombic
FeTCs FeS,
hexagonal cubic
hexagonal
a = 8.799-9.188 , = 4.50 4.60 varies with c o m p o s i t i o n a = 4.70-4.95 c = 7.70-8.15 varies with composition a = 4.751, c = 25.28 a = 10.910, c = 19.354 a = 5.099, b = 6.744 c = 4.5255 a = 6.882, c = 4.540
lS-Cr
[12]
:~- M n Fe~P (Mn + Cr)s 9{Ni + Feh~, 4Si~,.7 (Crt~,FesMozJC6, {Fe.Cr)e~C,,. (Cr,vFe4.sMo, s)C6, Cr_,~C6, (Cr,Fe.Mo),~C6 M 18C, F e ~ , C r l e M o m , {Fe,Nih6Cr,sMo4 FeCr. F e M o , F e C r M o . CrNiMo. CrFeMoNi
[ 10]
[5] [5]
Fe2Mo. Fes,,CrsSi s
[1.5]
FCTMO,, FeseMn,,Mo~e. Mnv,~MosSig9 Fe ~C
[91 [9]
FeTC~ FeS2
P
[10] [7J [5]
Mn¢,Si,
[9] [9] [9]
P
MTC 3 MrC
hexagonal fcc
MnS
fcc
a = 13.98. c = 4 . 5 2 5 a = 10.85-11.75 varies with composition tt = 5.224
M 6 C and M23C 6 carbides and oxide inclusions. Laves phase is probably the phase with an orange colour on the grain boundaries. M7C 3 carbides, on the other hand, precipitated in the matrix and on the grain boundaries. In the coloured surface photographs shown in Fig. 4(c), spherical carbides are seen distributed in the matrix of the ferrite phase. It is considered these precipitated phases are probably M 6 C metal carbides existing between 540 and 950°C in duplex alloys [5]. The nodular phases on the grain boundaries in Fig. 5(a) are probably M7C3 carbides. In Fig. 6, ,Z and G phases are seen besides the o" phase, inclusions and grainboundary carbides. The G phase is formed on the grain boundaries and in the ferrite matrix, whilst the Z phase is formed on the grain boundaries in the austenite phase. The G phases in the ferrite matrix are larger than those on the austenite-ferrite grain boundaries and are located on the dislocations in the ferrite matrix. Vitek et al. also reported that G phase, which is rich in Ni, Mn and Si, would form along the dislocations in ferrite and sometimes precipitate at the intersection between austenite and ferrite [7]. The dislocation etch pits are seen with difficulty as a triangular form in the ferrite matrix. In a body-centered cubic metal, dislocation etch
CrTC~ Fe~Mo~C
[5,1 l] [5]
MnS
[9]
pits in a (111) plane will appear as triangular-based pyramids with {110}-faced sides [8]. In the photographs in Fig. 4(d) and Fig. 5(b), inversion oil was not used, the blue-coloured phases in both figures being a phase. In Fig. 5(b), the red-coloured phases in band form are probably sulfur compounds. 3.2.
X-ray analysis
X-Ray diffractograms of specimens S I - $ 4 are seen in Fig. 7, from which diffractograms it has been understood that the heat treatments applied to the specimens did not change the crystal structure and the types of phases. However, the appearance or disappearance of the peaks numbered 3 - 4 and 6-7 may indicate that the phases with low percentage changed by a small amount. The 20 angles corresponding to the peak numbers in Fig. 5 are given in Table 3. There are small changes in the values of 20 angles, which means that the heat treatments caused changes in the compositions of the phases (especially of carbide phases). The changes in the chemical composition of the phases also change the lattice parameters by a small amount. Therefore, the heat treatments can be said to have caused a three-di-
244
M. Ceylan et al./Journal of Materials Processing Technology 69 (1997) 238-246
i,,,
,,,rrrl"ln,~,
I |, | ,,, :, iFi,
I .......
"1
.......
T'li'~''
'' '' 1
'
7
i
~
"
°Oj
[mW] -22
| -3.0
100
1t0
lO0
~0
300
~
400
450
[Ci
500
A •••ii•••••i••••i••••ii7i•i•••••iii•i••••iw•iii•••iii•iiii•i•iiiiiwi|•iiiiii•i•iii•i•iiv| [mW] - 2.0
i
\. li.go
-2.8
14.111,J/Ip
1
ul.lo
m
a
i itlllililiil
tl00
ill
iil|l
IllO
I I lie Iil
II00
I lili
li|lliiilll'il
190
itllll
380
ill|ll
_i lie
8S0
i llilili
iilai
400
| i I ilil
49¢
(C)
ill
I
S0O
B
Fig. 8. For legend please see facing page.
mensional lattice deformation in the phases, thus the mean values of 20 were used in the calculations. The common dependence on these angles of the common phases determined in metallographic examinations and the planes showing the precipitated phases, are given in Table 4, from which it can be seen for all of the specimens S I - $ 4 ~ ferrite and ), austenite are dominant phases. On the other hand, probable other phases and carbides are listed in Table 5 with crystal structures and lattice parameters.
3.3. DSC and DTA analysis In Fig. 8(a-d), DSC curves of the specimens in the same heating regime (10°C min-~) are shown. The behavior of the DSC curves of S1 and $3 are similar to each other (see Fig. 8(a) and (b)), whereas in SI only there are two small peaks at 462.5 and 469.7°C, and in $3 there exists only one peak. These peaks shows a phase formed of which the stoichiometry compound type existing at a constant temperature is very small.
M. Ceylan et al./ Journal of Materials Processing Technology 69 (1997) 238-246
245
[mW] - 2.5
l.O,~O
- 3.6
(C)
[mwJ - 3.o
.~
3. ~
. lode'@
- 3.3 ~| a nlannl aO0
nail ~60
e lalllaaalillJnaeAnll :SO0
a ! ~
on,el 400
4~
(C)
500
Fig. 8. DSC curves of specimens: (a) $1; (b) $3; (c) $2; and (d) $4; respectively.
DSC curves of the specimens $2 and $4 are shown in Fig. 8(c) and (d), from the comparison of which curves it can be concluded that a phase transformation occurred in $2 at 430.6°C, but that an athermal phase dissolved at 309.4°C in addition to the crystallization of two stoichiometry compounds at 393.7 and 418.9°C in specimen $4. This endothermic peak at 309.4°C may correspond to the dissolution of the martensite phase. When the DSC curves of S1 and $3, and $2 and $4 are compared, it is seen that S1 and $3 oxidized more during the annealing in the DSC instrument and that the precipitated phases were distributed over a very
large area. During the annealing, $2 and $4 did not oxidize because they were subjected to a similar treatment in DSC, but a large amount of precipitate phases formed at about 430°C. Carbides and the other phases absorb energy from the system when they dissolve and decrease the system's energy. The atoms of the dissolved phases diffuse to and ,~ phases, during which process they continue to take energy from the system. The temperatures at which the homogenizing is completed were determined from the DTA curves in Fig. 9 for S1, $2, $3 and $4 as 1123.7, 988.6, 959.4 and 928.6°C, respectively. Austen-
M. Ceylan et al./Journal of Materials Processing Technology 69 (1997) 238-246
246
Table 6 The results of the hardness measurements of tile specimens
-~.....,.
125
-~..."
||
q
dl 125
II 250
II 375
'in 500
• 625
/ 750
, 875
, 1000
• 1125
| 12=0
Specimens
S1
$2
$3
$4
Hardness (HV 5 kg)
251
212
257
239
caused by decrease in the amount of the carbon and chromium contents of the grains because of the formation of carbides, which are effective in the formation of martensite. Therefore, the hardness of specimen $3 does not show a great difference from that of SI.
r~]
Fig. 9. DTA curves of specimens SI-$4.
ite begins to give energy to the system during ~ + ? ~ transformation because of the nucleation of the new phase on the grain boundary, the temperatures for the greatest percentages being 1182.5, 1213.0, 1202.9 and 1219.2°C for specimens SI, $2, $3 and $4, respectively. In specimen S1, it has been determined that the ), to o~ transformation occurred in a small temperature gap and at a lower temperature compared to the other specimens, the reason for this being the lesser amount of the precipitated carbide in the grain than in the grain boundaries. In the annealed specimens ($2, $4) the peak widths are also greater as well as the peaks being of greater height. The reason for this is having allowed the amount of the precipitates in the grains to increase and the dissolving of the phases being difficult during the annealing. 3.4. Hardness measurements
As can be seen from Table 6, the greatest hardness values were measured in specimens S1 and $3, there being no considerable difference between these hardness values (251 HV for S1 and 257 HV for $3). The hardness of $2 is the lowest (212 HV), being the result of the distribution of some of the grain boundary carbides in the phases. Specimen $4 had a greater hardness when cooled in liquid nitrogan (239 HV), which may be due to the transformation of some of the austenite to martensite and increase in the internal tension. In the DSC curve of the same specimen in Fig. 6(d) the athermal phase melting at 309°C show the existence of martensite - - thought to exist in S4--in specimen $3 cooled in the liquid nitrogen, which can be
4. Conclusions Insufficient heat treatment causes the formation of the precipitate phases. The corrosion resistance, impact resistance and impact transition temperature are effected by these phases to a significant extent, therefore the period and the temperatures of the heat treatment have a great importance in obtaining the desired results.
References [1] G. Krauss, Steels: Heat Treatment and Processing Principles, ASM International, 1993. [2] M.R.M. Afshord, Corrosion-Resistance of AISI 316 Stainless Steel Castings, PhD Thesis, Dept. of Metallurgy, University of Brunel, 1983, pp. 12-14. [3] ] Charles, The Duplex Stainless Steels: Materials to Meet Your Needs, Beaune, France, 1991, pp. 3-48. [4] R.A. Lula, Stainless Steel, Fifth Printing, American Society for Metals, Metals Park, Ohio, 1993, pp. 89-93. [5] K. Mills et al. (Eds.), ASM Handbook: Metallography and Microstructures, Vol. 9, Formerly Ninth Edition, ASM International, 1992. [6] H.D. Solomon, Age hardening in a duplex stainless steel, in: R.A. Lula (Ed.), Duplex Stainless Steels, American Society for Metals, Metals Park, OH, 1983, pp. 12-15. [7] J.M. Vitek, S.A. David et al., Acta Metall. 39 (1991) 503-516. [8] G.F.V. Voort, Metallography: Principles and Practice, McGrawHill, New York, 1984. [9] V. Raghavan, Phase Diagrams of Ternary Iron Alloys, Part 2, The Indian Institute of Metals, Calcutta, 1988. [10] V. Raghavan, Phase Diagrams of Ternary Iron Alloys, Part 3, The Indian Institute of Metals, Calcutta, 1988. [11] G.V. Raynor, V.G. Rivlin, Phase Equilibria in Iron Ternary Alloys, The Institute of Metals, 1988. [12] C.S. Barrett, T.B. Massalski, Structure of Metals, Pergamon Press, Oxford, 1982.
