- Email: [email protected]
An assessment of carburization-decarburizatton behaviour of Fe-9Cr-Mo steels in a sodium environment
Journal of Nuclear Materials 110 (1982) l-10
North-Holland
Publishing Company
AN ASSESSMENT OF CARBURIZATION-DECARBURIZATrON MO STEELS IN A SODIUM ENVIRONMENT * A.
BEHAVIOUR
OF Fe-9Cr-
SALTELLI **, O.K. CHOPRA, and K. NATESAN
Materials
Science Division, Argonne National Luboratov,
Received 4 November
1981; accepted
Argonne, IL 60439, USA
1 April 1982
A critical assessment is made of the carburization-decarburization kinetics of Fe-9Cr-Mo steels exposed to a sodium environment, using the available information on carbide phase morphology, chromium activity in a ferrite matrix, chromium carbide activity in mixed carbides, carbon solubility in Cr-Mo ferritic steels, and activity-concentration relationships based on a-phase/M,C or Ma& carbide equilibrium. Experimental data are presented on the decarburization of Fe-SCr-Mo and Fe-9Cr-Mo steels at 973 K in a sodium environment to ascertain the long-term behaviour of these steels. The analysis shows that the decarburization of ferritic steels is largely dependent on the chemical reaction at the carbide/a interface and that at carbon activities <0.04, the rate is predominantly determined by the dissolution of (Fe, Mo),C carbides.
1. Introduction
Ferritic steels with 2-12 wt% chromium and l-2 wtX molybdenum contents are being considered for construction of sodium-heated steam generators and as alternative structural materials for liquid-sodium heattransport sytems, i.e. in liquid-metal fast-breeder reactors, fusion reactors, and central receiver solar-thermal power systems. The higher-chromium ferritic steels, in general, provide a greater resistance to carbon transfer than Fe-2$Zr-1 MO steel [1,2] and possess adequate mechanical properties. The carburization-decarburization kinetics of several commercial and high-purity Fe-9Cr- 1 to 2SMo ferritic steels (see compositions in table 1) were evaluated after exposure to a flowing sodium environment at temperatures between 773 and 973 K [3]. Carbon activity-concentration relationships for the high-purity and commercial Fe-9Cr-Mo steels were determined by equilibration of thin (- 70 pm) foils of the steels in flowing sodium and in CH,/H, gas mixtures of known carbon activities and subsequent analysis of the foils for total carbon using a combustion technique. A summary of the results obtained from these experiments is shown in
* Work supported by the US Department of Energy. **On a visiting appointment from Comitato Nazionale L’Energia Nucleare, Italy.
0022-3115/82/0000-0000/$02.75
Per
0 1982 North-Holland
figs. 1 and 2 for the high-purity and commercial Fe9Cr-Mo steels, respectively. In addition, kinetic experiments were conducted by exposure of 6 mm diameter rod samples of the steels at several temperatures and for time periods in the range 10.8-26.1 Ms. The exposed samples were machined, turnings were analyzed for total carbon, and carbon-concentration profiles were established as a function of sodium exposure time and temperature. Fig. 3 shows the carbon diffusivity values for the steels obtained from the kinetic experiments. The results show that the Fe-9Cr-Mo steels are much more resistant to carburization-decarburization in a sodium environment than the Fe-2$Cr1Mo steel. The carbon activity-concentration relationship and the effective diffusion coefficient for carbon (figs. l-3) can be used to examine the extent of carbon transfer in a secondary system of an LMFBR. However, the following observations suggest that a better understanding of the carburization-decarburization behaviour for the steels is required before data can be extrapolated to longer time periods which are of interest in modelling carbon transfer in secondary heat-transport systems: (a) the insensitivity of carbon concentration in the steels to variations in carbon activity in sodium in the range 0.05-0.001, (b) a plateau carbon concentration at low carbon activities which is dependent on initial carbon level in the steel (namely - 900 ppm for the high-purity alloys and -400 ppm for the HCM9M steel), and (c) the non-classical diffusion profiles for carbon con-
A. Salt&i et al. / Carburization
2
of Fe - 90 -MO steels
Table 1 Composition (wt%) of ferritic steels Steel
c
Cr
MO
Ni
Si
Mn
0.091 0.096
8.63 8.83
1.41 2.38
0.04 0.05
0.13 0.01
O.OtJ7 0.001
0.005 0.005
0.005 0.002
0.097 0.043 0.087 0.120 0.090
9.22 8.77 0.58 8.48 9.46
1.00 2.12 0.80 2.26 2.13
0.08 0.08 0.07
0.08 0.30 0.17 0.30 0.33
0.38 0.53 0.44 1.16 0.88
0.15 0.11 0.51 0.39
0.22 0.15 0.22 0.18
V
‘Nb
Other
High-purity
Fe-9 Cr-1.5 MO Fe-9 (Z-2.5 MO Commerciai
Heat 91887 HCM9M ESRXA 3177 ESRXA 3089 EM-12
0.038 N 0.46 W, 0.04 Ti, 0.066 N 0.042 N 0.026 N
centration in the kinetic experiments. In the present paper, a thermodynamic analysis of the carburization-decarburization processes in Fe9Cr-MO steels was made to establish the validity of the experimental carbon activity-concentration relationships. Additional information is presented on the decarburization of Fe-SCr-Mo and Fe-9Cr-Mo steels at 973 K to ascertain the long-term behaviour of the steels.
CARBON ACTIVITY
Fig. 1. Carbon acti~ty-~n~ntration relationship for higb-purity Fe-9Cr-Mo steels at temperatures between 973 and 773 K.
/ Fe-SOCr-MOSTEEL
I
(bl
TEMPERATURE ("C) 600 5x) I I
650 I
l ESRXA 3089
1
l EM-12
1
0 ESRXA 3177
500 I
DECARBURIZED
0
I.1 CARBONACTIVITY
Fig. 2. Carbon activity-concentration relationship for ESRXA 3089, EM-12, and HCM9M steels at temperatures between 973 and 773 K.
1.2 103/T(K)
1.3
Fig. 3. Effective carbon diffusivities calculated from the carburization-decarburization rate constants for Fe-9Cr-Mo steels.
A. Sake& et al. / Curbarizationof Fe - 90 -MO steels
2. Thermodynamic
analysis
The relevant information required in the analysis of the carburization-decarburization processes in Fe-CrMO ferritic steels includes the following: (a) carbidephase morphology, (b) chro~um activity in a ferritic matrix, (c) chromium carbide activity in mixed (Fe, Cr),,C, carbide, (d) carbon solubility in Cr-Mo ferritic steels, and (e) phase diagram and activity-concentration relationships based on a! phase-M,C or M,,C, carbide equilibrium. Each of these areas will be examined in detail in the subsequent sections. 2. I. Carbide-phase morphoiogy According to studies reported by Andrews et al. [4], both M,C and h&C, carbides are stable in Fe-9Cr-Mo steels and the relative amount of each will depend on the MO/C atomic ratio in the steel. Kuo [5] reported that M,C and M,,C, coexist in Fe-MO-C systems when the MO/C atom ratio is between 1.5 and 3 and M,C is the only stable carbide phase at ratios above 3. The results of Kuo were obtained on low (up to 1.16 wt%) chromium alioys; the M,,C, carbide phase in this chromium concentration range has the composition Mo,Fe,,C,. In a Cr-Mo steel with a higher chromium content, the composition of M&Z, phase can vary over a range established by Mo,Fe,,C,-and Cr,,C, [6]. In Fe-9Cr-Mo steels, carbides of the composition Cr,,Fe,C, have been established by the chemical analysis of carbides extracted from thermally aged samples. Molybdenum is observed only in carbides of the M,C type; also, the maximum contribution of Cr to the M,C carbide is a Cr/metal ratio of 0.3, corresponding to a composition Cr,,s(Fe, MO)& for the carbide [6]. It has been fairly well established that M,C phase forms as an intermediate carbide in the thermal aging of Fe-9CrMO steels [3]. As in the case of Fe-ZaCr-1Mo steel, the decarburization resistance of an Fe-9Cr-Mo alloy is significantly larger if the material has a stable carbide structure (achieved by the tempering treatment). If material with an intermediate carbide phase (M,C) is exposed to a sodium en~ronment, breakdown of M,C and decarburization of the steel can occur, leading to accelerated formation of M,C phase. While significant data are available on carbon acivities associated with M,,C, carbide with a wide variation in metal atom content (Cr, Fe, MO, etc.) almost no thermodynamic info~ation is available for the M,C phase. As a result, the carbon activity-concentration relationships are evaluated based on M,,C, phase alone and subsa quently M,C carbide is used in the calculations to
3
account for the difference between the experimentally observed and theoretically computed carbon activityconcentration relationships. 2.2. Chromium activity in Fe-Cr-Mo
alloys
Very few experimental data are available on the variation in chromium activity with chromium concentration in ferritic steels. Therefore, we attempt here to evaluate the chromium activity coefficient for an Fe-9Cr steel using the available high-temperature thermodynamic information. For the solid-solution reaction crcs)
16OOK + Crt8-iron)
atX,,=O.l,
(1)
we have from [7], AH, = 4536 cal mol-’ AS, = 6.381 cal mol-’ K-t, where AH, and AS, are the enthalpy and entropy change for reaction (1). Similarly, to calculate AH and AS at lower temperatures, we can write for the reaction Crl;, -) Cr$e K
(2)
the following enthalpy and entropy expressions, obtained by interpolation of data selected from Hultgren [8] in the vicinity of T= 600 K: AH2 = 8202 - 6.87( T - 600) A&=7.79-O.Ol(T-600). Interpolating the data of Kendall and Hultgren [9] in a similar manner for the reaction involving chromium dissolution in S-iron at 1600 K and o-iron at a temperature T, we can write Cr(YEj( Xc, = 0. I > --, Cr& iron)
(Xc, = 0.1);
(3)
for this case, AH,=
AS,=
-9474+8.32(T-600)
-9.45+0.0119(T-600).
Summing the enthalpy and entropy equations from the above expressions, we obtain the following overall reaction for the dissolution of chromium in a-iron: Cr& -+ Cr(T iron)
(Xc,=O.l).
(4)
Now
1.45 T-(3.58+0.0019 T) X T. (5) Since AG = RTln( $rXc,), the expression for the chromium activity coefficient becomes AG=2394+
1204.8 In y& = 1.23 + 7 - 9.56 X lop4 T
(6)
A. Saltelli et al. / Carburization of Fe - 9Cr - MO steels
4 Table 2 Chromium
activity
coefficient
in a-phase TEMPERATURE 973K A CARBON ACTIVITY
Temperature
ATWCARBIDE EQ",L,BR,L,M
(K)
1.0
4.65 5.86 7.72 10.7
973 873 773 673
for a steel with Xc, = 0.1 (i.e. Cr = 9 wt%). Values of -& are listed in table2 for several temperatures of interest. Similar values for y& could be obtained by directly extrapolating the AC data of [7] for reaction (1) to lower temperatures, assuming AGT=AH1600K-AS’600K(Z--
2.3. Chromium carbide activity carbide
1600).
(7)
in the mixed (Fe, Cr)C, 10-44 0
2
4
6
B
IO
The (Fe, Cr),,C, carbide has been observed as either a stable or metastable carbide phase together with M,C phase in many of the Fe-Cr-Mo steels [3]. The average composition of the mixed carbide is Fe,Cr,,C,. In order to generate a carbon activity-concentration relationship for the ferritic steel, it is essential to establish the carbide composition and activity in equilibrium with the a phase. Even though a significant spread exists in the data for the free energy of formation of ternary (Cr, Fe) carbides [lo-121, which has recently been discussed by Benz et al. [ 121, an attempt can be made to calculate the activity of binary Cr,,C, carbide in the ternary mixed carbide using the available information. The ideal mixing model for the ternary carbides proposed by Richardson [13], and successfully applied by Butler et al. [ 141 and Natesan and Kassner [ 151, was used for the evaluation of Cr,,C, activity in the ternary carbide. According to this model, the Cr,,C, activity in a mixed carbide of composition Cr,,Fe,C, is given by
concentrations at the solubility limit in these alloys. Darken and Ryba [16] used a thermogravimetric technique to obtain carbon activity-concentration curves by equilibrating Fe-Cr-Mo (up to 3 wt% Cr) steel specimens with several hydrogen-methane gas mixtures. The carbon activity and concentration values corresponding to the onset of carbide precipitation at 973 K are shown in fig. 4 as a function of chromium content of the alloys. Although no details are reported on the morphology of the precipitated carbides, the data were extrapolated to higher chromium levels of interest in the present work. The expression (based on extrapolation) at 973 K are
ac,,,c,
In C,sa’(ppm)
= @s/23)23.
(8)
CHROMIUM CONCENTRATION(wl%)
Fig. 4. Carbon saturation activity and concentration as a function of chromium concentration in Fe-Cr-C alloys at 973 K.
= 5.30 - 0.438 Cr(wt%).
2.5. Carbon activity-concentration
(9)
relationships
2.4. Carbon solubility in Cr-Mo ferritic steels Few experimental data are available on the solubility of carbon in Cr-Mo ferritic steels, owing to the complex morphology of the carbide phase(s) and low carbon
The purpose of this section is to use the above information to establish the carbon activity-concentration relationship in a-phase in equilibrium with M,,C, carbide at 973 K. The value for free energy of formation
A. Saltelli et al. / Carburization
concentration in the a-phase. A similar relationship between total carbon and carbon activity can be obtained using eq. (12). As an example, the calculated values of carbon and chromium concentrations in aphase and the carbide compositions are listed in table 4 for a total carbon concentration of 0.6 wt% in the steel. The total chromium concentrations listed in the table were calculated by the stoichiometric relationship (in weight units) Crrot = Cr, + 13.0 X (C,, - C,), assuming a composition of Cr,,Fe,C, for the carbide. The carbide phase compositions were calculated from the above data using the relations
Table 3 Thermodynamic properties of Cr,,C, AS @l/gatom K)
AG at 973 K @al/g atom C)
Reference
- 16380
1.54
- 17900
-10800a' -114OOb’ -12800 -11750
3.05 -
-12300c' - 12900= - 15800 - 13250”
P71 I181 1191 1201 (191
AH (Cal/g atom C)
a) Revised by Dawson and Sale [ 191.
CrJwt%)
b, Work of Alekseer quoted in [ 191. c, Computed using entropy data of Kelley et al. [ 171.
C) = - 12630 + 2.30 X T(K). (10)
For the reaction qCr
+ C = iCr,,C,,
the equilibrium
(11)
constant
K can be written
as
4$,
=--216at973K. 23/6 acrac
= x,C,(wtW)
Cr,,(wt%)
of Cr,,C, obtained by several investigators are listed in table 3. Since there is no basis for selecting the results of a particular investigation, the enthalpy and entropy data were averaged to obtain the expression AGcr2,c6(cal/g-atom
5
of Fe - 90 -MO steels
(12)
With the use of eqs. (6) and (8) to calculate ~6~ and uc123c6, respectively, we can obtain a relationship between total carbon concentration and carbon in the u-phase (eq. (9)) for a steel with a given chromium
= &Cr,(wt%)
+ (1 - X,)C,(wt%) + (1 - X,)Cr,(wt%).
(13) (14)
The data of table4 enable construction of the phase diagram shown in fig. 5. For a given Cr concentration in the alloy, the tie lines of fig. 5 establish a relation between total carbon concentration and carbon activity. A section of the diagram corresponding to a total chromium content of 9 wt% is shown in fig. 6; the dashed portion of the diagram is calculated according to Henry’s law, i.e. a, = (a~‘/C~)C,. The shaded area of fig. 6 encompasses the experimental data obtained at 973 K. The analysis shows a fairly good agreement between the calculated and experimentally determined carbon concentration values at carbon acivities greater than - 0.04. At carbon activities less than 0.04, the significant difference between the calculated and experimental results observed in the figure cannot be accounted for by composition variation of the M,,C, carbide. As mentioned earlier, M,C carbide phase predominates as a stable phase in Fe-Cr-Mo ferritic steels when decarburized in a liquid-sodium environment [3,12]. Although the M,C formula is generally used
Table 4 Calculated values of a-phase and carbide composition for a total carbon content
of 0.6 wt% in Fe-Cr-Mo
ferritic
steel
C&pm)
Cr,(wt%)
Ch
a,
x,
Cr,
m
129 83 54 35 22 15 9.3 6.0 3.9 1.6
1 2 3 4 5 6 7 8 9 11
8.63 9.69 10.73 11.75 12.77 13.78 14.79 15.19 16.79 18.80
1.0 1.0 1.0 0.89 0.38 0.19 0.10 0.062 0.040 0.018
0.1058 0.1066 0.1071 0.1074 0.1076 0.1077 0.1078 0.1078 0.1079 0.1080
73.1 74.0 75.2 76.2 77.2 78.3 79.3 80.3 81.2 83.3
18.1 18.3 18.6 18.8 19.1 19.3 19.5 19.8 20.0 20.5
6
A. SalteNi et ai. / Curburization of Fe-9Cr-MO 100
2 '
60-
p
IO 30 50 _ 90 ENLARGED&-CARBID? (Cr,iFe,Mol23-,,,)CG
L.__..-1
I
2
3
4
5
6
CARBON CONCENTRATION lwt%l
Fig. 5. Fe-O-C phase diagram at 973 K. The tie lines in the a-carbide region show the compositions of the a and carbide phases for sweral carbon activity values in the system.
together with the designation of 9 carbides, the actual chemical composition of this carbide has been the subject of several investigations. The first extensive study
CARBON ACTIVITY
Fig. 6. Carbon concentration in Fe-9Cr-Mo of carbon activity at 973 K. The shaded
steel as a function area in the figure
encompasses all the experimental data obtained at ANL.
steels
on the ? carbides was conducted by Kuo [5,22]. He sintered double carbides of the Group IV transition metals (designated as A) with metals of Groups V and VI ‘(designated as B) at temperatures in the range 1773 to 2073 K and observed mixed-carbide compositions of the type A,B,C and A,B,C, which were designated 7, and na carbides, respectively. The (MO, Fe) mixed carbide of composition Fe,Mo.,C observed in Cr-Mo ferritic steels corresponds to q2-type carbide. By aging Fe-MO steels with up to 8 wt% MO content at 973 K for time periods in the range 180 s to 5000 h, Kuo [S] also observed Fe-Mo 9, carbides with compositions ranging between Mo,Fe,C and Mo,Fe,C in the final product. In his experiments, M,C was observed to form only when the MO/C ratio exceeded 1.2, and was generally accompanied by M,,C,. The n carbide existed alone only for a MO/C ratio W3.3. The M,,C, carbide was detected as an intermediate phase, but it always dissolved in his experiments with longer exposures, resulting in M,C or MoC. In a study of molybdenum carbide transformation kinetics conducted by aging 3.5 wt% MO steels at 873 and 973 K for times ranging from 1 to 1000 h, Ridal and Quarrel (231 observed q,(M,C) carbide as the product of transformation of MO& phase that was initially present. The authors suggested that the transformation of M,C to M,C occurs in situ and that a MO/C ratio of at least 2 is required for complete transformation. They also found that the theoretical yield of n carbide based on the composition Fe,Mo$ was much larger than the experimental yield, while the chemical analyses of the carbides indicated an iron/metal atom ratio of -0.5. From these results and some considerations about M,C crystal structure, the authors concluded that a formula Fe2,4Mo,,,C better represented the actual carbide, corresponding to M, C ,.a5 rather than to the M,C formula. The spread in the composition of the n carbides has been emphasized by Fraker and Stadelmayer [24], who attempted to establish the stability of the n phase in several systems. Their phase diagram for the Fe-MO-C system at 1273 K indicates that Mo,Fe,C (i.e. 11,) is the only stable carbide. The carbide phase Mo,Fe,C was not observed in their studies, which is in disagreement with the results of Kuo [22]. The composition of the n phase in Fe-2iCr-1Mo steel after low-temperature aging and after exposure to a sodium environment (decarburizing condition) was evaluated by Klueh and Leitnaker [21] and Leitnaker et al. [25]. In these studies, specimens of Fe-2fCr-1Mo steel were both decarburized in sodium and thermally aged in a helium atmosphere at 839 K for 26500 h. The carbides in the initial material (prior to exposure) were
A. Sahel/i et al. / Carburization
M,C, M,C, and M&J,, with M,,C, constituting a major fraction. In the thermally aged specimens, the extracted carbide phase contained 60% M,C and 40% M,,C,. The decarburized specimens exhibited some differences in the carbide composition between those in the decarburized layer and in the interior of the specimens. In the decarburized zone, as much as 94% of the carbides were of the M,C type and the balance were M&Z,. In the interior of the specimen, the M,C content ranged between 70 and 86%. Decarburization not only favoured the formation of M,C in the carbon-depleted region, but also decreased the volume fraction of carbide precipitates when compared with the thermally aged specimens. Based on the results obtained in longterm aging experiments, the authors concluded that possibly the end point of Baker and Nutting’s [26] carbide transformation sequence is never reached and that for all practical purpose the material always contains M,,C, together with M,C. Furthermore, they reported a value approaching 4 rather than 6 for the metal to carbon ratio in the 7 phase of Fe-2aCr-1Mo steel after exposure at 839 K. This observation is in the direction of the composition M,,,C for the 7 carbide reported by Ridal and Quarrel [23]. Based on the above results, three conclusions can be drawn regarding carbide stability in ferritic steel: (a) The MO/C atomic ratio in the q carbide varies with exposure temperature; the MO content in the carbide is generally larger at higher temperatures. It is not clear at present whether q, and q, carbide structures are two separate compounds or represent a wide variation in composition from a stoichiometric formula. (b) As the MO/C ratio decreases in the alloy, the metal to carbon ratio in the 9 carbide also decreases. M,C and M,,C, will coexist even after long aging times unless the MO/C ratio is > 3.3. (c) M,C and MC phases are responsible for the decarburization of Fe-24 Cr- 1MO and Fe-9Cr-Mo steels. At low carbon activities, M,C is more stable than M,,C,; therefore, under decarburizing conditions in a sodium environment, the ferritic steel exhibited accelerated formation of M,C when compared with thermally aged material. With the information presented thus far on the role of MO on carbide precipitation in ferritic steels, we can examine the carbon activity-concentration relationship in the steels, especially at carbon activities less than 0.04. Since no information is avialable on the activity coefficient for MO in these steels or on the free energy of formation of M, C phase, one can write an expression for equilibrium between a-phase and M,C carbide as (MO&Z,
= k,
(15)
7
of Fe -9Cr -MO steels
1.0:
,
, ,
(
,
,
,
,
I
I
I
lllll
TOTAL CA
3
-
I
Illlll~
1
I
t000
0.1
0.01
CARBON ACTIVITY
Fig. 7. Total carbon concentration values in Fe-9Cr-2.5Mo steel as a function of carbon activity. Also shown are the curves calculated on the basis of equilibrium between a phase and M,C or M,,C,. where (MO) is the concentration of MO in the a-phase and the constant k includes the MsC activity in the (Fe, Cr, Mo),C mixed carbide. Referring to fig. 7, in which the experimental data for Fe-9Cr-2SMo steel are reported, we can assume a carbide of the M,C type at a low carbon activity value of 0.002 (the MO/C ratio in the steel corresponding to this activity is -4.0). Assuming a carbide composition of MO, Fe,C and using a total carbon value of 0.07 wt% at a carbon activity of 0.002 (see fig. 7), one obtains a calculated value of 1.38 wt% for molybdenum concentration in the a-phase. Insertion of this value for (MO) in eq. (15) results in a k value of 0.0139. Now we have for the a-phase/(Fe, Cr),,C, carbide system (using eqs. (12), (8), and (6)), (Cr)23’6a,
= 180
and for the a-phase/(Fe, (MO)&
(16) Mo),C
carbide
= 0.0139,
system, (17)
where (Ci) and (MO) aqe the chromium and molybdenum concentrations in a-phase. If a and b represent the carbon concentraiions (in wt%) in M,,C, and M,C, respectively, a = [9 - (Cr)]
X
(12 X 6/52 X 18)
and b= [2.5 - (MO)] X(12/95,94X2). Calculations concentration 0.002 to 0.4. and are also fixed MO/C
(18)
were made to obtain the carbon activityrelationship over a wide activity range of The calculated values are listed in table5 plotted in fig. 7. Despite the assumption of and Cr/C ratios in the two carbides, the
A. Saltelli et al. / Carburiration of Fe - 9Cr - MO steels
8
table6. Prior to the exposure, the alloys were solution annealed for 1 h at 1323 K, air cooled, tempered for 1 h at 1923 K, and air cooled. The purpose of these tests is to ascertain whether or not an equilibrium carbide configuration is attained in the alloys after long-time exposure at 973 K and to compare the carbon concentration values in the exposed specimens with the activity-concentration relationship established earlier. Samples of the four alloys were exposed at 973 K to flowing sodium in a stainless steel loop, the details of which are given elsewhere [3]. The carbon concentration in the sodium was - 0.05 ppm, which corresponds to a at 973 K. The results, carbon activity of -0.0014 namely the total carbon concentration (Cr.,,) and the change in the carbon concentration (Cinit’a’ - Crol) of the alloys after different exposure times, are plotted against the square root of time in fig. 8. The behaviour of each alloy with respect to decarburization and carbide morphology can now be examined. After the initial tempering, all four alloys contained M,,C, as a dominant carbide with some MO& and M,C. During short-time exposures in a sodium environment, the MO& carbide dissolves, leading to decarburization of all the alloys. However, during this time period, the transformation M,,C, -+ M,C occurs in those steels in which the MO/C ratio is sufficiently high, i.e. in Fe-9Cr-2.5Mo and Fe-5Cr-2.5Mo steels. The formation of M,C phase slows the decarburization process in these steels. Furthermore, a higher chromium content in the steel results in a larger volume fraction of M,,C, carbide (during tempering) which also seems to slow the decarburization process in a sodium environment. A comparison of the results for Fe-9Cr-0.5Mo and Fe-SCr-0.5Mo steels also indicates a lower decarburization rate in steels with a higher chromium content. The results also show that steels with a higher molybdenum content reach a stable carbide structure (M6C) after fairly short exposure times and further loss of carbon from the steel is almost negligible. On the other hand, for steels with lower MO contents (e.g. Fe-9Cr-O.SMo), the rate of decarburization is
Table 5 Calculated values for a-phase composition and total carbon as a function of carbon activity Concentration (wt%) ac
(Cr)
(Mo)
a
b
C l-0,
0.002 0.004 0.01 0.02 0.04 0.07 0.10 0.2 0.4
9.00 9.00 9.00 9.00 8.97 7.76 7.07 5.90 4.92
1.38 1.23 1.06 0.94 0.84 0.76 0.72 0.64 0.57
0 0 0 0 0.002 0.096 0.149 0.239 0.314
0.070 0.079 0.090 0.098 0.104 0.109 0.111 0.116 0.121
0.07 0.079 0.090 0.098 0.106 0.204 0.260 0.355 0.434
calculations show a good agreement with the experimental data on the carbon activity-concentration relationship. The results show that at low carbon activities the MO-based carbides will dictate the carburization/ decarburization process while the Cr-based carbides will dominate at high carbon activities. For steels with lower molybdenum contents, the actual carbide phase in the low activity range is likely to be composed of M,C and M,,C, rather than M,C alone. Nevertheless, the analysis can be used effectively to establish the carbon activity-concentration relationships for the Fe-Cr-Mo ferritic steels.
3. Experimental results 3.1. Decarburization
experiments
Four different high-purity Fe-Cr-Mo steels (75 pm thick foil specimens) were exposed to a decarburizing environment at 973 K for time periods between 6 and 1000 h. The compositions of the alloys are listed in
Table 6 Composition steel
(wtl)
of steels used in the experimental C
Cr
MO
program Ni
Si
Mn
Nb
V
Fe
0.001 0.002 o.cXl2 0.005
0.005 0.005 0.005 0.005
0.002 0.001 0.001 0.002
Balance Balance Balance Balance
Fe-9Cr-2.5Mo
0.096
8.83
2.38
0.05
0.01
Fe-9Cr-0.5Mo Fe-SCr-2Mo Fe-SCr-0.5Mo
0.098 0.092 0.094
9.46 4.32 4.48
0.50 1.87 0.46
0.07 0.04 0.06
0.01 0.09 0.13
A. SakIli et al. / Carburizationof Fe - 9Cr -MO steefs EXPOSURE IO
50
TIME.1 thl
loo
500
I
I
loo0
~~ 0
02
04
06
0.8
I.0
I Ma 5 MO
*Fe-50-0.5 + Fe-5Cr-2 a c__o,-.
1200
1.2
1.4
1.6
113
2.0
Ji (103s)
EXPOSURE
L 1200
1000
TIME.1 Ih)
IO
50
ml
500
1
I
I
I
0 0 0 A
1000
Fe-5 Cr-0.5 MO Fe-5Cr-2.5Mo Fe-9 0-0.5 MO Fe-90-2 5Mo
9
the carbide particles [27~29]. According to the Hillert [27] model, the order of magnitude of the dissolutionprecipitation time can be estimated by the time required for the complete equilibration of chromium within the grain. The grain size (h) was G 10 pm in all four steels used in the experimental programme. Using the Cr diffusion coefficient value of 3.2 X lo-l3 cm’ s-’ at 973 K [30], the time for complete redist~bution of chromium in the grain is t = (h/2)‘/4 L& = 55 h. The molybdenum diffusivity is slightly higher than that of chromium [30], which results in a lower redistribution time in the case of MO-based carbides. Carbon diffusion within the foils can be neglected in view of much larger values for the interstitial carbon diffusivity in ferritic steels. The time dependence of the carbon concentration values (fig. 8) shows that the dissolution rate of carbides is dependent more on the chemical reaction at the carbide/a-phase interface than on the subsititutional element diffusion in the a-phase. The lower rates of decarbu~zation can be attributed to the much smaller thermodynamic gradient for carbon at the carbide/a interface. This is in contrast with the observations on austenitic stainless steels in which the decarburization process and the carbide dissolution are determined by the chromium diffusion in the austenite away from the carbide particle.
4. summary 02
0.4
0.6
0.6
I.0
1.2
1.4
1.6
1.8
2.0
Fig. 8. Time dependence of total carbon concentration (top) and mass of carbon loss (bottom) for several Fe-Cr-Mo steels exposed to a flowing sodium environment at 973 K.
determined by the rate off dissolution of M,,C, and the carbon loss continues even after a 1000 h exposure. The decarburization rate constant of K, = 1.26 X IO-* g cms2 sell2 is in good agreement with the earlier results [3]. The Fe-SCr-OSMo steel decarburizes much faster, owing to the lower volume fraction of the M,& phase and relatively faster dissolution of the carbide particles. 3.2. Models for decarburization kinetics Models for the carbide pr~ipitation-evolution in austenitic stainless steels have been developed which are based upon chromium diffusion toward or away from
A thermodynamic analysis of the carburization-decarburization processes in Fe-9Cr-Mo ferritic steels is presented to establish the validity of expe~men~lly determined carbon activity-concentration relationships. Experimental data are presented on the decarburization of Fe-SCr-Mo and Fe-9Cr-Mo steels at 973 K to ascertain the long-term behaviour of these steels in a sodium environment. Based on the detailed analyses, it can be concluded that the curve of fig. 1 represents the carbon activity-concentration relationship for normalized and tempered Fe-9Cr-2SMo steel. The decarburization rate constants evaluated from data obtained after - 500 h exposures at 973 K can be used effectively to describe the carbon loss from the material, especially in the design of components with long service life. The minimum exposure time for the evaluation of decarburization rate constants will increase for lower test temperatures, but the information presented in [4] and the curves in Figs. I and 2 can be used to describe the decarburization behaviour of intermediate chromiummolybdenum ferritic steels exposed to a sodium environment.
A. Saltelli et al. / Carburization
10
a,
ap
-
cm pat cc C Tot Cr, Cr Tot Cr,
-
Di
-
AGi AH,
-
h
-
K
-
KD m
_
AS, T t xc x, Vi
activity of an element
or a phase i. concentration. carbon concentration in a-phase, wt%. saturation concentration of carbon in cx phase,
_ carbon activity at saturation
ppm. carbon concentration in carbide, wt%. total carbon concentration in steel, wt% chromium concentration in a phase, wt%. total chromium concentration in steel, wt%. chromium concentration in carbide, wt%. diffusion coefficient for element i, cm2 s-‘. free-energy change for reaction i, cal mol-‘. enthalpy change for reaction i, cal mol-‘. grain size. thermodynamic equilibrium constant. decarburization rate constant, g cmp2 s-‘12. stoichiometric coefficient in the formula
CrmPe(23--mJC6. - entropy change for reaction i, cal mol-’ K. - absolute temperature, _ time, s. - mole fraction of carbide in steel. - mole fraction of element i. - activity coefficient of element i.
K-‘.
Acknowledgements The fellowship of the Comitato Nazionale Per L‘Energia Nucleare (Italian Nuclear Authority) to one of the authors (AS) is gratefully acknowledged. D.L. Kink assisted with experimental program. The work was supported by the US Department of Energy.
References [I] O.K. Chopra, K. Natesan, and T.F. Kassner, in: Proc. Int. Conf. on Liquid Metal Technology in Energy Production, CONF-760503-P2 (1976) p. 730. [2] G. Menken, E.D. Grosser, and E. Te Hessen, in: Proc. Int. Conf. on Ferritic Steels for Fast Reactor Steam Generators (BNES, London, 1978) p. 264. [3] O.K. Chopra, K. Natesan, and T.F. Kassner, J. Nucl. Mater. 96 (1981) 269.
of Fe - 90 -MO steels [41K.W. Andrews, H. Hughes, and D.J. Dyson, J. Iron Steel Inst. 210 (1972) 337. [51 K. Kuo, J. Iron Steel Inst. 173 (1953) 363. 161J.H. Woodhead and A.G. Quarrel, J. Iron Steel Inst. 203 (1965) 605. [71R. Hultgren et al., Selected Values of the Thermodynamic Properties of Binary Alloys (American Society for Metals, Metals Park, OH, 1973). PI R. Hultgren et al., Selected Values of the Thermodynamic Properties of the Elements (American Society for Metals, Metals Park, OH, 1973). 191W.B. Kendall and R. Hultgren, Trans. Am. Sot. Met. 53 (1961) 207. Russ. Metall. Part 1 1101V.I. Alekseer and L.A. Shvartsman, (1965) 117. [IllH. Tuma, P. Grobner and K. Lobl, Arch. Eisenhiittenwes. 40 (1969) 727. 1121R.Benz, J.F. Elliott, and J. Chipman, Met. Trans. 5 (1974) 2235. J. Iron Steel Inst. 175 (1953) 33. 1131F.D. Richardson, 1141J.F. Butler, C.L. McCabe, and H.W. Paxton, Trans. AIME 221 (1961) 479. [I51K. Natesan and T.F. Kassner, J. Nucl. Mater. 37 (1970) 223. Pennsylvania [I61 L.S. Darken and E.R. Ryba, EPRI-AF-1179, State University (1979). 1171 K.K. Kelley, F.S. Boericke, G.E. Moore, E.H. Huffman, and W.M. Bangert, US Bureau of Mines Techn. Paper 662 (1949). iI81A.D. Mah, US Bureau of Mines R. I. 7217 (1969). 1191W.M. Dawson and F.R. Sale, Met. Trans. 8A (1977) 15. PO1 A.D. Kulkarni and W.L. Worrell, Met. Trans. 3 (1972) 2363. Met. Trans. 6A (1975) 1211R.L. Klueh and J.M. Leitnaker, 2089. WI K. Kuo, Acta Metall. 1 (1953) 301. ~231 K.A. Ridal and A.G. Quarrel, J. Iron Steel Inst. 200 (1962) 359. and H.H. Stadelmayer, Trans. AIME 245 ~241A.C.Fraker (1969) 847. 1251J.M. Leitnaker, R.L. Klueh, and W.R. Laing, Met. Trans. 6A (1975) 1949. WI R.G. Baker and J. Nutting, J. Iron Steel Inst: 192 (1959) 257. 1271 C. Stawstrom and M. Hillert, J. Iron Steel Inst. 207 (1969) 77. [28] R. Gullberg, J. Iron Steel Inst. 211 (1973) 59. [29] M. Hillert, K. Nilsson, and L.E. Tbmdahl, J. Iron Steel Inst. 209 (1971) 49. [30] P.J. Alberry and C.W. Haworth, Met. Sci. 8 (1974) 407.
North-Holland
Publishing Company
AN ASSESSMENT OF CARBURIZATION-DECARBURIZATrON MO STEELS IN A SODIUM ENVIRONMENT * A.
BEHAVIOUR
OF Fe-9Cr-
SALTELLI **, O.K. CHOPRA, and K. NATESAN
Materials
Science Division, Argonne National Luboratov,
Received 4 November
1981; accepted
Argonne, IL 60439, USA
1 April 1982
A critical assessment is made of the carburization-decarburization kinetics of Fe-9Cr-Mo steels exposed to a sodium environment, using the available information on carbide phase morphology, chromium activity in a ferrite matrix, chromium carbide activity in mixed carbides, carbon solubility in Cr-Mo ferritic steels, and activity-concentration relationships based on a-phase/M,C or Ma& carbide equilibrium. Experimental data are presented on the decarburization of Fe-SCr-Mo and Fe-9Cr-Mo steels at 973 K in a sodium environment to ascertain the long-term behaviour of these steels. The analysis shows that the decarburization of ferritic steels is largely dependent on the chemical reaction at the carbide/a interface and that at carbon activities <0.04, the rate is predominantly determined by the dissolution of (Fe, Mo),C carbides.
1. Introduction
Ferritic steels with 2-12 wt% chromium and l-2 wtX molybdenum contents are being considered for construction of sodium-heated steam generators and as alternative structural materials for liquid-sodium heattransport sytems, i.e. in liquid-metal fast-breeder reactors, fusion reactors, and central receiver solar-thermal power systems. The higher-chromium ferritic steels, in general, provide a greater resistance to carbon transfer than Fe-2$Zr-1 MO steel [1,2] and possess adequate mechanical properties. The carburization-decarburization kinetics of several commercial and high-purity Fe-9Cr- 1 to 2SMo ferritic steels (see compositions in table 1) were evaluated after exposure to a flowing sodium environment at temperatures between 773 and 973 K [3]. Carbon activity-concentration relationships for the high-purity and commercial Fe-9Cr-Mo steels were determined by equilibration of thin (- 70 pm) foils of the steels in flowing sodium and in CH,/H, gas mixtures of known carbon activities and subsequent analysis of the foils for total carbon using a combustion technique. A summary of the results obtained from these experiments is shown in
* Work supported by the US Department of Energy. **On a visiting appointment from Comitato Nazionale L’Energia Nucleare, Italy.
0022-3115/82/0000-0000/$02.75
Per
0 1982 North-Holland
figs. 1 and 2 for the high-purity and commercial Fe9Cr-Mo steels, respectively. In addition, kinetic experiments were conducted by exposure of 6 mm diameter rod samples of the steels at several temperatures and for time periods in the range 10.8-26.1 Ms. The exposed samples were machined, turnings were analyzed for total carbon, and carbon-concentration profiles were established as a function of sodium exposure time and temperature. Fig. 3 shows the carbon diffusivity values for the steels obtained from the kinetic experiments. The results show that the Fe-9Cr-Mo steels are much more resistant to carburization-decarburization in a sodium environment than the Fe-2$Cr1Mo steel. The carbon activity-concentration relationship and the effective diffusion coefficient for carbon (figs. l-3) can be used to examine the extent of carbon transfer in a secondary system of an LMFBR. However, the following observations suggest that a better understanding of the carburization-decarburization behaviour for the steels is required before data can be extrapolated to longer time periods which are of interest in modelling carbon transfer in secondary heat-transport systems: (a) the insensitivity of carbon concentration in the steels to variations in carbon activity in sodium in the range 0.05-0.001, (b) a plateau carbon concentration at low carbon activities which is dependent on initial carbon level in the steel (namely - 900 ppm for the high-purity alloys and -400 ppm for the HCM9M steel), and (c) the non-classical diffusion profiles for carbon con-
A. Salt&i et al. / Carburization
2
of Fe - 90 -MO steels
Table 1 Composition (wt%) of ferritic steels Steel
c
Cr
MO
Ni
Si
Mn
0.091 0.096
8.63 8.83
1.41 2.38
0.04 0.05
0.13 0.01
O.OtJ7 0.001
0.005 0.005
0.005 0.002
0.097 0.043 0.087 0.120 0.090
9.22 8.77 0.58 8.48 9.46
1.00 2.12 0.80 2.26 2.13
0.08 0.08 0.07
0.08 0.30 0.17 0.30 0.33
0.38 0.53 0.44 1.16 0.88
0.15 0.11 0.51 0.39
0.22 0.15 0.22 0.18
V
‘Nb
Other
High-purity
Fe-9 Cr-1.5 MO Fe-9 (Z-2.5 MO Commerciai
Heat 91887 HCM9M ESRXA 3177 ESRXA 3089 EM-12
0.038 N 0.46 W, 0.04 Ti, 0.066 N 0.042 N 0.026 N
centration in the kinetic experiments. In the present paper, a thermodynamic analysis of the carburization-decarburization processes in Fe9Cr-MO steels was made to establish the validity of the experimental carbon activity-concentration relationships. Additional information is presented on the decarburization of Fe-SCr-Mo and Fe-9Cr-Mo steels at 973 K to ascertain the long-term behaviour of the steels.
CARBON ACTIVITY
Fig. 1. Carbon acti~ty-~n~ntration relationship for higb-purity Fe-9Cr-Mo steels at temperatures between 973 and 773 K.
/ Fe-SOCr-MOSTEEL
I
(bl
TEMPERATURE ("C) 600 5x) I I
650 I
l ESRXA 3089
1
l EM-12
1
0 ESRXA 3177
500 I
DECARBURIZED
0
I.1 CARBONACTIVITY
Fig. 2. Carbon activity-concentration relationship for ESRXA 3089, EM-12, and HCM9M steels at temperatures between 973 and 773 K.
1.2 103/T(K)
1.3
Fig. 3. Effective carbon diffusivities calculated from the carburization-decarburization rate constants for Fe-9Cr-Mo steels.
A. Sake& et al. / Curbarizationof Fe - 90 -MO steels
2. Thermodynamic
analysis
The relevant information required in the analysis of the carburization-decarburization processes in Fe-CrMO ferritic steels includes the following: (a) carbidephase morphology, (b) chro~um activity in a ferritic matrix, (c) chromium carbide activity in mixed (Fe, Cr),,C, carbide, (d) carbon solubility in Cr-Mo ferritic steels, and (e) phase diagram and activity-concentration relationships based on a! phase-M,C or M,,C, carbide equilibrium. Each of these areas will be examined in detail in the subsequent sections. 2. I. Carbide-phase morphoiogy According to studies reported by Andrews et al. [4], both M,C and h&C, carbides are stable in Fe-9Cr-Mo steels and the relative amount of each will depend on the MO/C atomic ratio in the steel. Kuo [5] reported that M,C and M,,C, coexist in Fe-MO-C systems when the MO/C atom ratio is between 1.5 and 3 and M,C is the only stable carbide phase at ratios above 3. The results of Kuo were obtained on low (up to 1.16 wt%) chromium alioys; the M,,C, carbide phase in this chromium concentration range has the composition Mo,Fe,,C,. In a Cr-Mo steel with a higher chromium content, the composition of M&Z, phase can vary over a range established by Mo,Fe,,C,-and Cr,,C, [6]. In Fe-9Cr-Mo steels, carbides of the composition Cr,,Fe,C, have been established by the chemical analysis of carbides extracted from thermally aged samples. Molybdenum is observed only in carbides of the M,C type; also, the maximum contribution of Cr to the M,C carbide is a Cr/metal ratio of 0.3, corresponding to a composition Cr,,s(Fe, MO)& for the carbide [6]. It has been fairly well established that M,C phase forms as an intermediate carbide in the thermal aging of Fe-9CrMO steels [3]. As in the case of Fe-ZaCr-1Mo steel, the decarburization resistance of an Fe-9Cr-Mo alloy is significantly larger if the material has a stable carbide structure (achieved by the tempering treatment). If material with an intermediate carbide phase (M,C) is exposed to a sodium en~ronment, breakdown of M,C and decarburization of the steel can occur, leading to accelerated formation of M,C phase. While significant data are available on carbon acivities associated with M,,C, carbide with a wide variation in metal atom content (Cr, Fe, MO, etc.) almost no thermodynamic info~ation is available for the M,C phase. As a result, the carbon activity-concentration relationships are evaluated based on M,,C, phase alone and subsa quently M,C carbide is used in the calculations to
3
account for the difference between the experimentally observed and theoretically computed carbon activityconcentration relationships. 2.2. Chromium activity in Fe-Cr-Mo
alloys
Very few experimental data are available on the variation in chromium activity with chromium concentration in ferritic steels. Therefore, we attempt here to evaluate the chromium activity coefficient for an Fe-9Cr steel using the available high-temperature thermodynamic information. For the solid-solution reaction crcs)
16OOK + Crt8-iron)
atX,,=O.l,
(1)
we have from [7], AH, = 4536 cal mol-’ AS, = 6.381 cal mol-’ K-t, where AH, and AS, are the enthalpy and entropy change for reaction (1). Similarly, to calculate AH and AS at lower temperatures, we can write for the reaction Crl;, -) Cr$e K
(2)
the following enthalpy and entropy expressions, obtained by interpolation of data selected from Hultgren [8] in the vicinity of T= 600 K: AH2 = 8202 - 6.87( T - 600) A&=7.79-O.Ol(T-600). Interpolating the data of Kendall and Hultgren [9] in a similar manner for the reaction involving chromium dissolution in S-iron at 1600 K and o-iron at a temperature T, we can write Cr(YEj( Xc, = 0. I > --, Cr& iron)
(Xc, = 0.1);
(3)
for this case, AH,=
AS,=
-9474+8.32(T-600)
-9.45+0.0119(T-600).
Summing the enthalpy and entropy equations from the above expressions, we obtain the following overall reaction for the dissolution of chromium in a-iron: Cr& -+ Cr(T iron)
(Xc,=O.l).
(4)
Now
1.45 T-(3.58+0.0019 T) X T. (5) Since AG = RTln( $rXc,), the expression for the chromium activity coefficient becomes AG=2394+
1204.8 In y& = 1.23 + 7 - 9.56 X lop4 T
(6)
A. Saltelli et al. / Carburization of Fe - 9Cr - MO steels
4 Table 2 Chromium
activity
coefficient
in a-phase TEMPERATURE 973K A CARBON ACTIVITY
Temperature
ATWCARBIDE EQ",L,BR,L,M
(K)
1.0
4.65 5.86 7.72 10.7
973 873 773 673
for a steel with Xc, = 0.1 (i.e. Cr = 9 wt%). Values of -& are listed in table2 for several temperatures of interest. Similar values for y& could be obtained by directly extrapolating the AC data of [7] for reaction (1) to lower temperatures, assuming AGT=AH1600K-AS’600K(Z--
2.3. Chromium carbide activity carbide
1600).
(7)
in the mixed (Fe, Cr)C, 10-44 0
2
4
6
B
IO
The (Fe, Cr),,C, carbide has been observed as either a stable or metastable carbide phase together with M,C phase in many of the Fe-Cr-Mo steels [3]. The average composition of the mixed carbide is Fe,Cr,,C,. In order to generate a carbon activity-concentration relationship for the ferritic steel, it is essential to establish the carbide composition and activity in equilibrium with the a phase. Even though a significant spread exists in the data for the free energy of formation of ternary (Cr, Fe) carbides [lo-121, which has recently been discussed by Benz et al. [ 121, an attempt can be made to calculate the activity of binary Cr,,C, carbide in the ternary mixed carbide using the available information. The ideal mixing model for the ternary carbides proposed by Richardson [13], and successfully applied by Butler et al. [ 141 and Natesan and Kassner [ 151, was used for the evaluation of Cr,,C, activity in the ternary carbide. According to this model, the Cr,,C, activity in a mixed carbide of composition Cr,,Fe,C, is given by
concentrations at the solubility limit in these alloys. Darken and Ryba [16] used a thermogravimetric technique to obtain carbon activity-concentration curves by equilibrating Fe-Cr-Mo (up to 3 wt% Cr) steel specimens with several hydrogen-methane gas mixtures. The carbon activity and concentration values corresponding to the onset of carbide precipitation at 973 K are shown in fig. 4 as a function of chromium content of the alloys. Although no details are reported on the morphology of the precipitated carbides, the data were extrapolated to higher chromium levels of interest in the present work. The expression (based on extrapolation) at 973 K are
ac,,,c,
In C,sa’(ppm)
= @s/23)23.
(8)
CHROMIUM CONCENTRATION(wl%)
Fig. 4. Carbon saturation activity and concentration as a function of chromium concentration in Fe-Cr-C alloys at 973 K.
= 5.30 - 0.438 Cr(wt%).
2.5. Carbon activity-concentration
(9)
relationships
2.4. Carbon solubility in Cr-Mo ferritic steels Few experimental data are available on the solubility of carbon in Cr-Mo ferritic steels, owing to the complex morphology of the carbide phase(s) and low carbon
The purpose of this section is to use the above information to establish the carbon activity-concentration relationship in a-phase in equilibrium with M,,C, carbide at 973 K. The value for free energy of formation
A. Saltelli et al. / Carburization
concentration in the a-phase. A similar relationship between total carbon and carbon activity can be obtained using eq. (12). As an example, the calculated values of carbon and chromium concentrations in aphase and the carbide compositions are listed in table 4 for a total carbon concentration of 0.6 wt% in the steel. The total chromium concentrations listed in the table were calculated by the stoichiometric relationship (in weight units) Crrot = Cr, + 13.0 X (C,, - C,), assuming a composition of Cr,,Fe,C, for the carbide. The carbide phase compositions were calculated from the above data using the relations
Table 3 Thermodynamic properties of Cr,,C, AS @l/gatom K)
AG at 973 K @al/g atom C)
Reference
- 16380
1.54
- 17900
-10800a' -114OOb’ -12800 -11750
3.05 -
-12300c' - 12900= - 15800 - 13250”
P71 I181 1191 1201 (191
AH (Cal/g atom C)
a) Revised by Dawson and Sale [ 191.
CrJwt%)
b, Work of Alekseer quoted in [ 191. c, Computed using entropy data of Kelley et al. [ 171.
C) = - 12630 + 2.30 X T(K). (10)
For the reaction qCr
+ C = iCr,,C,,
the equilibrium
(11)
constant
K can be written
as
4$,
=--216at973K. 23/6 acrac
= x,C,(wtW)
Cr,,(wt%)
of Cr,,C, obtained by several investigators are listed in table 3. Since there is no basis for selecting the results of a particular investigation, the enthalpy and entropy data were averaged to obtain the expression AGcr2,c6(cal/g-atom
5
of Fe - 90 -MO steels
(12)
With the use of eqs. (6) and (8) to calculate ~6~ and uc123c6, respectively, we can obtain a relationship between total carbon concentration and carbon in the u-phase (eq. (9)) for a steel with a given chromium
= &Cr,(wt%)
+ (1 - X,)C,(wt%) + (1 - X,)Cr,(wt%).
(13) (14)
The data of table4 enable construction of the phase diagram shown in fig. 5. For a given Cr concentration in the alloy, the tie lines of fig. 5 establish a relation between total carbon concentration and carbon activity. A section of the diagram corresponding to a total chromium content of 9 wt% is shown in fig. 6; the dashed portion of the diagram is calculated according to Henry’s law, i.e. a, = (a~‘/C~)C,. The shaded area of fig. 6 encompasses the experimental data obtained at 973 K. The analysis shows a fairly good agreement between the calculated and experimentally determined carbon concentration values at carbon acivities greater than - 0.04. At carbon activities less than 0.04, the significant difference between the calculated and experimental results observed in the figure cannot be accounted for by composition variation of the M,,C, carbide. As mentioned earlier, M,C carbide phase predominates as a stable phase in Fe-Cr-Mo ferritic steels when decarburized in a liquid-sodium environment [3,12]. Although the M,C formula is generally used
Table 4 Calculated values of a-phase and carbide composition for a total carbon content
of 0.6 wt% in Fe-Cr-Mo
ferritic
steel
C&pm)
Cr,(wt%)
Ch
a,
x,
Cr,
m
129 83 54 35 22 15 9.3 6.0 3.9 1.6
1 2 3 4 5 6 7 8 9 11
8.63 9.69 10.73 11.75 12.77 13.78 14.79 15.19 16.79 18.80
1.0 1.0 1.0 0.89 0.38 0.19 0.10 0.062 0.040 0.018
0.1058 0.1066 0.1071 0.1074 0.1076 0.1077 0.1078 0.1078 0.1079 0.1080
73.1 74.0 75.2 76.2 77.2 78.3 79.3 80.3 81.2 83.3
18.1 18.3 18.6 18.8 19.1 19.3 19.5 19.8 20.0 20.5
6
A. SalteNi et ai. / Curburization of Fe-9Cr-MO 100
2 '
60-
p
IO 30 50 _ 90 ENLARGED&-CARBID? (Cr,iFe,Mol23-,,,)CG
L.__..-1
I
2
3
4
5
6
CARBON CONCENTRATION lwt%l
Fig. 5. Fe-O-C phase diagram at 973 K. The tie lines in the a-carbide region show the compositions of the a and carbide phases for sweral carbon activity values in the system.
together with the designation of 9 carbides, the actual chemical composition of this carbide has been the subject of several investigations. The first extensive study
CARBON ACTIVITY
Fig. 6. Carbon concentration in Fe-9Cr-Mo of carbon activity at 973 K. The shaded
steel as a function area in the figure
encompasses all the experimental data obtained at ANL.
steels
on the ? carbides was conducted by Kuo [5,22]. He sintered double carbides of the Group IV transition metals (designated as A) with metals of Groups V and VI ‘(designated as B) at temperatures in the range 1773 to 2073 K and observed mixed-carbide compositions of the type A,B,C and A,B,C, which were designated 7, and na carbides, respectively. The (MO, Fe) mixed carbide of composition Fe,Mo.,C observed in Cr-Mo ferritic steels corresponds to q2-type carbide. By aging Fe-MO steels with up to 8 wt% MO content at 973 K for time periods in the range 180 s to 5000 h, Kuo [S] also observed Fe-Mo 9, carbides with compositions ranging between Mo,Fe,C and Mo,Fe,C in the final product. In his experiments, M,C was observed to form only when the MO/C ratio exceeded 1.2, and was generally accompanied by M,,C,. The n carbide existed alone only for a MO/C ratio W3.3. The M,,C, carbide was detected as an intermediate phase, but it always dissolved in his experiments with longer exposures, resulting in M,C or MoC. In a study of molybdenum carbide transformation kinetics conducted by aging 3.5 wt% MO steels at 873 and 973 K for times ranging from 1 to 1000 h, Ridal and Quarrel (231 observed q,(M,C) carbide as the product of transformation of MO& phase that was initially present. The authors suggested that the transformation of M,C to M,C occurs in situ and that a MO/C ratio of at least 2 is required for complete transformation. They also found that the theoretical yield of n carbide based on the composition Fe,Mo$ was much larger than the experimental yield, while the chemical analyses of the carbides indicated an iron/metal atom ratio of -0.5. From these results and some considerations about M,C crystal structure, the authors concluded that a formula Fe2,4Mo,,,C better represented the actual carbide, corresponding to M, C ,.a5 rather than to the M,C formula. The spread in the composition of the n carbides has been emphasized by Fraker and Stadelmayer [24], who attempted to establish the stability of the n phase in several systems. Their phase diagram for the Fe-MO-C system at 1273 K indicates that Mo,Fe,C (i.e. 11,) is the only stable carbide. The carbide phase Mo,Fe,C was not observed in their studies, which is in disagreement with the results of Kuo [22]. The composition of the n phase in Fe-2iCr-1Mo steel after low-temperature aging and after exposure to a sodium environment (decarburizing condition) was evaluated by Klueh and Leitnaker [21] and Leitnaker et al. [25]. In these studies, specimens of Fe-2fCr-1Mo steel were both decarburized in sodium and thermally aged in a helium atmosphere at 839 K for 26500 h. The carbides in the initial material (prior to exposure) were
A. Sahel/i et al. / Carburization
M,C, M,C, and M&J,, with M,,C, constituting a major fraction. In the thermally aged specimens, the extracted carbide phase contained 60% M,C and 40% M,,C,. The decarburized specimens exhibited some differences in the carbide composition between those in the decarburized layer and in the interior of the specimens. In the decarburized zone, as much as 94% of the carbides were of the M,C type and the balance were M&Z,. In the interior of the specimen, the M,C content ranged between 70 and 86%. Decarburization not only favoured the formation of M,C in the carbon-depleted region, but also decreased the volume fraction of carbide precipitates when compared with the thermally aged specimens. Based on the results obtained in longterm aging experiments, the authors concluded that possibly the end point of Baker and Nutting’s [26] carbide transformation sequence is never reached and that for all practical purpose the material always contains M,,C, together with M,C. Furthermore, they reported a value approaching 4 rather than 6 for the metal to carbon ratio in the 7 phase of Fe-2aCr-1Mo steel after exposure at 839 K. This observation is in the direction of the composition M,,,C for the 7 carbide reported by Ridal and Quarrel [23]. Based on the above results, three conclusions can be drawn regarding carbide stability in ferritic steel: (a) The MO/C atomic ratio in the q carbide varies with exposure temperature; the MO content in the carbide is generally larger at higher temperatures. It is not clear at present whether q, and q, carbide structures are two separate compounds or represent a wide variation in composition from a stoichiometric formula. (b) As the MO/C ratio decreases in the alloy, the metal to carbon ratio in the 9 carbide also decreases. M,C and M,,C, will coexist even after long aging times unless the MO/C ratio is > 3.3. (c) M,C and MC phases are responsible for the decarburization of Fe-24 Cr- 1MO and Fe-9Cr-Mo steels. At low carbon activities, M,C is more stable than M,,C,; therefore, under decarburizing conditions in a sodium environment, the ferritic steel exhibited accelerated formation of M,C when compared with thermally aged material. With the information presented thus far on the role of MO on carbide precipitation in ferritic steels, we can examine the carbon activity-concentration relationship in the steels, especially at carbon activities less than 0.04. Since no information is avialable on the activity coefficient for MO in these steels or on the free energy of formation of M, C phase, one can write an expression for equilibrium between a-phase and M,C carbide as (MO&Z,
= k,
(15)
7
of Fe -9Cr -MO steels
1.0:
,
, ,
(
,
,
,
,
I
I
I
lllll
TOTAL CA
3
-
I
Illlll~
1
I
t000
0.1
0.01
CARBON ACTIVITY
Fig. 7. Total carbon concentration values in Fe-9Cr-2.5Mo steel as a function of carbon activity. Also shown are the curves calculated on the basis of equilibrium between a phase and M,C or M,,C,. where (MO) is the concentration of MO in the a-phase and the constant k includes the MsC activity in the (Fe, Cr, Mo),C mixed carbide. Referring to fig. 7, in which the experimental data for Fe-9Cr-2SMo steel are reported, we can assume a carbide of the M,C type at a low carbon activity value of 0.002 (the MO/C ratio in the steel corresponding to this activity is -4.0). Assuming a carbide composition of MO, Fe,C and using a total carbon value of 0.07 wt% at a carbon activity of 0.002 (see fig. 7), one obtains a calculated value of 1.38 wt% for molybdenum concentration in the a-phase. Insertion of this value for (MO) in eq. (15) results in a k value of 0.0139. Now we have for the a-phase/(Fe, Cr),,C, carbide system (using eqs. (12), (8), and (6)), (Cr)23’6a,
= 180
and for the a-phase/(Fe, (MO)&
(16) Mo),C
carbide
= 0.0139,
system, (17)
where (Ci) and (MO) aqe the chromium and molybdenum concentrations in a-phase. If a and b represent the carbon concentraiions (in wt%) in M,,C, and M,C, respectively, a = [9 - (Cr)]
X
(12 X 6/52 X 18)
and b= [2.5 - (MO)] X(12/95,94X2). Calculations concentration 0.002 to 0.4. and are also fixed MO/C
(18)
were made to obtain the carbon activityrelationship over a wide activity range of The calculated values are listed in table5 plotted in fig. 7. Despite the assumption of and Cr/C ratios in the two carbides, the
A. Saltelli et al. / Carburiration of Fe - 9Cr - MO steels
8
table6. Prior to the exposure, the alloys were solution annealed for 1 h at 1323 K, air cooled, tempered for 1 h at 1923 K, and air cooled. The purpose of these tests is to ascertain whether or not an equilibrium carbide configuration is attained in the alloys after long-time exposure at 973 K and to compare the carbon concentration values in the exposed specimens with the activity-concentration relationship established earlier. Samples of the four alloys were exposed at 973 K to flowing sodium in a stainless steel loop, the details of which are given elsewhere [3]. The carbon concentration in the sodium was - 0.05 ppm, which corresponds to a at 973 K. The results, carbon activity of -0.0014 namely the total carbon concentration (Cr.,,) and the change in the carbon concentration (Cinit’a’ - Crol) of the alloys after different exposure times, are plotted against the square root of time in fig. 8. The behaviour of each alloy with respect to decarburization and carbide morphology can now be examined. After the initial tempering, all four alloys contained M,,C, as a dominant carbide with some MO& and M,C. During short-time exposures in a sodium environment, the MO& carbide dissolves, leading to decarburization of all the alloys. However, during this time period, the transformation M,,C, -+ M,C occurs in those steels in which the MO/C ratio is sufficiently high, i.e. in Fe-9Cr-2.5Mo and Fe-5Cr-2.5Mo steels. The formation of M,C phase slows the decarburization process in these steels. Furthermore, a higher chromium content in the steel results in a larger volume fraction of M,,C, carbide (during tempering) which also seems to slow the decarburization process in a sodium environment. A comparison of the results for Fe-9Cr-0.5Mo and Fe-SCr-0.5Mo steels also indicates a lower decarburization rate in steels with a higher chromium content. The results also show that steels with a higher molybdenum content reach a stable carbide structure (M6C) after fairly short exposure times and further loss of carbon from the steel is almost negligible. On the other hand, for steels with lower MO contents (e.g. Fe-9Cr-O.SMo), the rate of decarburization is
Table 5 Calculated values for a-phase composition and total carbon as a function of carbon activity Concentration (wt%) ac
(Cr)
(Mo)
a
b
C l-0,
0.002 0.004 0.01 0.02 0.04 0.07 0.10 0.2 0.4
9.00 9.00 9.00 9.00 8.97 7.76 7.07 5.90 4.92
1.38 1.23 1.06 0.94 0.84 0.76 0.72 0.64 0.57
0 0 0 0 0.002 0.096 0.149 0.239 0.314
0.070 0.079 0.090 0.098 0.104 0.109 0.111 0.116 0.121
0.07 0.079 0.090 0.098 0.106 0.204 0.260 0.355 0.434
calculations show a good agreement with the experimental data on the carbon activity-concentration relationship. The results show that at low carbon activities the MO-based carbides will dictate the carburization/ decarburization process while the Cr-based carbides will dominate at high carbon activities. For steels with lower molybdenum contents, the actual carbide phase in the low activity range is likely to be composed of M,C and M,,C, rather than M,C alone. Nevertheless, the analysis can be used effectively to establish the carbon activity-concentration relationships for the Fe-Cr-Mo ferritic steels.
3. Experimental results 3.1. Decarburization
experiments
Four different high-purity Fe-Cr-Mo steels (75 pm thick foil specimens) were exposed to a decarburizing environment at 973 K for time periods between 6 and 1000 h. The compositions of the alloys are listed in
Table 6 Composition steel
(wtl)
of steels used in the experimental C
Cr
MO
program Ni
Si
Mn
Nb
V
Fe
0.001 0.002 o.cXl2 0.005
0.005 0.005 0.005 0.005
0.002 0.001 0.001 0.002
Balance Balance Balance Balance
Fe-9Cr-2.5Mo
0.096
8.83
2.38
0.05
0.01
Fe-9Cr-0.5Mo Fe-SCr-2Mo Fe-SCr-0.5Mo
0.098 0.092 0.094
9.46 4.32 4.48
0.50 1.87 0.46
0.07 0.04 0.06
0.01 0.09 0.13
A. SakIli et al. / Carburizationof Fe - 9Cr -MO steefs EXPOSURE IO
50
TIME.1 thl
loo
500
I
I
loo0
~~ 0
02
04
06
0.8
I.0
I Ma 5 MO
*Fe-50-0.5 + Fe-5Cr-2 a c__o,-.
1200
1.2
1.4
1.6
113
2.0
Ji (103s)
EXPOSURE
L 1200
1000
TIME.1 Ih)
IO
50
ml
500
1
I
I
I
0 0 0 A
1000
Fe-5 Cr-0.5 MO Fe-5Cr-2.5Mo Fe-9 0-0.5 MO Fe-90-2 5Mo
9
the carbide particles [27~29]. According to the Hillert [27] model, the order of magnitude of the dissolutionprecipitation time can be estimated by the time required for the complete equilibration of chromium within the grain. The grain size (h) was G 10 pm in all four steels used in the experimental programme. Using the Cr diffusion coefficient value of 3.2 X lo-l3 cm’ s-’ at 973 K [30], the time for complete redist~bution of chromium in the grain is t = (h/2)‘/4 L& = 55 h. The molybdenum diffusivity is slightly higher than that of chromium [30], which results in a lower redistribution time in the case of MO-based carbides. Carbon diffusion within the foils can be neglected in view of much larger values for the interstitial carbon diffusivity in ferritic steels. The time dependence of the carbon concentration values (fig. 8) shows that the dissolution rate of carbides is dependent more on the chemical reaction at the carbide/a-phase interface than on the subsititutional element diffusion in the a-phase. The lower rates of decarbu~zation can be attributed to the much smaller thermodynamic gradient for carbon at the carbide/a interface. This is in contrast with the observations on austenitic stainless steels in which the decarburization process and the carbide dissolution are determined by the chromium diffusion in the austenite away from the carbide particle.
4. summary 02
0.4
0.6
0.6
I.0
1.2
1.4
1.6
1.8
2.0
Fig. 8. Time dependence of total carbon concentration (top) and mass of carbon loss (bottom) for several Fe-Cr-Mo steels exposed to a flowing sodium environment at 973 K.
determined by the rate off dissolution of M,,C, and the carbon loss continues even after a 1000 h exposure. The decarburization rate constant of K, = 1.26 X IO-* g cms2 sell2 is in good agreement with the earlier results [3]. The Fe-SCr-OSMo steel decarburizes much faster, owing to the lower volume fraction of the M,& phase and relatively faster dissolution of the carbide particles. 3.2. Models for decarburization kinetics Models for the carbide pr~ipitation-evolution in austenitic stainless steels have been developed which are based upon chromium diffusion toward or away from
A thermodynamic analysis of the carburization-decarburization processes in Fe-9Cr-Mo ferritic steels is presented to establish the validity of expe~men~lly determined carbon activity-concentration relationships. Experimental data are presented on the decarburization of Fe-SCr-Mo and Fe-9Cr-Mo steels at 973 K to ascertain the long-term behaviour of these steels in a sodium environment. Based on the detailed analyses, it can be concluded that the curve of fig. 1 represents the carbon activity-concentration relationship for normalized and tempered Fe-9Cr-2SMo steel. The decarburization rate constants evaluated from data obtained after - 500 h exposures at 973 K can be used effectively to describe the carbon loss from the material, especially in the design of components with long service life. The minimum exposure time for the evaluation of decarburization rate constants will increase for lower test temperatures, but the information presented in [4] and the curves in Figs. I and 2 can be used to describe the decarburization behaviour of intermediate chromiummolybdenum ferritic steels exposed to a sodium environment.
A. Saltelli et al. / Carburization
10
a,
ap
-
cm pat cc C Tot Cr, Cr Tot Cr,
-
Di
-
AGi AH,
-
h
-
K
-
KD m
_
AS, T t xc x, Vi
activity of an element
or a phase i. concentration. carbon concentration in a-phase, wt%. saturation concentration of carbon in cx phase,
_ carbon activity at saturation
ppm. carbon concentration in carbide, wt%. total carbon concentration in steel, wt% chromium concentration in a phase, wt%. total chromium concentration in steel, wt%. chromium concentration in carbide, wt%. diffusion coefficient for element i, cm2 s-‘. free-energy change for reaction i, cal mol-‘. enthalpy change for reaction i, cal mol-‘. grain size. thermodynamic equilibrium constant. decarburization rate constant, g cmp2 s-‘12. stoichiometric coefficient in the formula
CrmPe(23--mJC6. - entropy change for reaction i, cal mol-’ K. - absolute temperature, _ time, s. - mole fraction of carbide in steel. - mole fraction of element i. - activity coefficient of element i.
K-‘.
Acknowledgements The fellowship of the Comitato Nazionale Per L‘Energia Nucleare (Italian Nuclear Authority) to one of the authors (AS) is gratefully acknowledged. D.L. Kink assisted with experimental program. The work was supported by the US Department of Energy.
References [I] O.K. Chopra, K. Natesan, and T.F. Kassner, in: Proc. Int. Conf. on Liquid Metal Technology in Energy Production, CONF-760503-P2 (1976) p. 730. [2] G. Menken, E.D. Grosser, and E. Te Hessen, in: Proc. Int. Conf. on Ferritic Steels for Fast Reactor Steam Generators (BNES, London, 1978) p. 264. [3] O.K. Chopra, K. Natesan, and T.F. Kassner, J. Nucl. Mater. 96 (1981) 269.
of Fe - 90 -MO steels [41K.W. Andrews, H. Hughes, and D.J. Dyson, J. Iron Steel Inst. 210 (1972) 337. [51 K. Kuo, J. Iron Steel Inst. 173 (1953) 363. 161J.H. Woodhead and A.G. Quarrel, J. Iron Steel Inst. 203 (1965) 605. [71R. Hultgren et al., Selected Values of the Thermodynamic Properties of Binary Alloys (American Society for Metals, Metals Park, OH, 1973). PI R. Hultgren et al., Selected Values of the Thermodynamic Properties of the Elements (American Society for Metals, Metals Park, OH, 1973). 191W.B. Kendall and R. Hultgren, Trans. Am. Sot. Met. 53 (1961) 207. Russ. Metall. Part 1 1101V.I. Alekseer and L.A. Shvartsman, (1965) 117. [IllH. Tuma, P. Grobner and K. Lobl, Arch. Eisenhiittenwes. 40 (1969) 727. 1121R.Benz, J.F. Elliott, and J. Chipman, Met. Trans. 5 (1974) 2235. J. Iron Steel Inst. 175 (1953) 33. 1131F.D. Richardson, 1141J.F. Butler, C.L. McCabe, and H.W. Paxton, Trans. AIME 221 (1961) 479. [I51K. Natesan and T.F. Kassner, J. Nucl. Mater. 37 (1970) 223. Pennsylvania [I61 L.S. Darken and E.R. Ryba, EPRI-AF-1179, State University (1979). 1171 K.K. Kelley, F.S. Boericke, G.E. Moore, E.H. Huffman, and W.M. Bangert, US Bureau of Mines Techn. Paper 662 (1949). iI81A.D. Mah, US Bureau of Mines R. I. 7217 (1969). 1191W.M. Dawson and F.R. Sale, Met. Trans. 8A (1977) 15. PO1 A.D. Kulkarni and W.L. Worrell, Met. Trans. 3 (1972) 2363. Met. Trans. 6A (1975) 1211R.L. Klueh and J.M. Leitnaker, 2089. WI K. Kuo, Acta Metall. 1 (1953) 301. ~231 K.A. Ridal and A.G. Quarrel, J. Iron Steel Inst. 200 (1962) 359. and H.H. Stadelmayer, Trans. AIME 245 ~241A.C.Fraker (1969) 847. 1251J.M. Leitnaker, R.L. Klueh, and W.R. Laing, Met. Trans. 6A (1975) 1949. WI R.G. Baker and J. Nutting, J. Iron Steel Inst: 192 (1959) 257. 1271 C. Stawstrom and M. Hillert, J. Iron Steel Inst. 207 (1969) 77. [28] R. Gullberg, J. Iron Steel Inst. 211 (1973) 59. [29] M. Hillert, K. Nilsson, and L.E. Tbmdahl, J. Iron Steel Inst. 209 (1971) 49. [30] P.J. Alberry and C.W. Haworth, Met. Sci. 8 (1974) 407.