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Tensile and microstructural behavior of solute-modified manganese-stabilized austenitic steels
Materials Science and Engineering, A127 (1990) 17-31
17
Tensile and Microstructural Behavior of Solute-modified Manganese-stabilized Austenitic Steels R. L. KLUEH and P. J. MAZIASZ Metals and Ceramics Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6376 (U.S.A.) (Received September 5, 1989; in revised form December 4, 1989)
Abstract
A program is under way to develop a nickelfree austenitic stainless steel for fusion reactor applications. Previous work showed that a stable austenite alloy is possible in a compositional regime of Fe-(20-25)Mn-(12-14)Cr-(0.1-0.25)C. As a base composition for further alloying for strength and irradiation resistance, an Fe-20Mn12Cr-0.25C composition was chosen. Tensile properties for this base composition were comparable to those of type 316 stainless steel. To improve strength and irradiation resistance, closely controlled quantities of titanium, tungsten, vanadium, phosphorus and boron were added to the base alloy. Various combinations of these alloying elements improved the tensile properties of the new steels relative to type 316 stainless steel in both the annealed and 20% cold-worked conditions. These solute combinations also produced MC and M23 C6 precipitation in some alloys during annealing at 1050 and 1150 °C.
I. Introduction
Elements that make up the alloys proposed for the first wall and blanket structure of a magnetic fusion reactor become activated by the high energy neutrons created by the fusion reaction in the plasma. These radioactive reactor components must be properly disposed of after their service lifetime. The complexity of the waste disposal procedure depends on the time required for the induced radioactivity to decay to levels that are no longer hazardous to people and the environment. The more rapidly the radioactivity decays, the simpler the disposal task becomes. Part 61 of Title 10 of the U.S. Code of Federal Regulations (10CFR61) was prepared by the Nuclear Regulatory Commission for radioactive 0921-5093/90/$3.50
fission product disposal, Using 10CFR61, it is possible to define compositional limits for alloys that can be disposed of b y relatively simple shallow land burial techniques, as opposed to the much more difficult and expensive deep geologic burial. The objective of the present work is to develop new alloys that meet guidelines for shallow land burial. Some common alloying elements that decay over a long period of time after neutron activation are nickel, molybdenum, nitrogen, copper and niobium. An alloy development program is in progress [1] to develop fast induced radioactivity decay (FIRD) austenitic stainless steels to replace the high nickel stainless steels that are among the present candidate alloys for fusion applications. To produce an FIRD austenitic steel, manganese was proposed as a replacement for nickel, although it was recognized that this may prove difficult, because manganese is not as strong an austenite stabilizer as nickel and because the metallurgical effects of nickel and manganese in the alloy may not otherwise be equivalent [1]. The first attempts to use manganese as a replacement for nickel in stainless steels were made in the 1930s [2]. Since manganese and nickel behave differently in austenitic steels, commercially available high manganese stainless steels contain some nickel and use high concentrations of nitrogen to produce austenite stability. Other groups around the world are also working to develop high manganese, reduced activation austenitic steels [3-7]. However, in much of that work nitrogen is used for austenite stabilization [3, 4, 6] and in other instances small amounts of nickel are allowed [3, 4]. In keeping with the effort to meet the criteria established by the 10CFR61 guidelines, neither nitrogen nor nickel was used for austenite stabilization in the present work. © Elsevier Sequoia/Printed in The Netherlands
18
In the strategy proposed for developing FIRD steels [1], the first task was a determination of F e - M n - C r - C compositions that would produce a fully austenitic microstructure [1]. Such a base composition could then be modified with minor element additions to obtain strength and irradiation resistance at least equal to that found in current candidate fusion reactor structural materials. By investigating the microstructures of a series of Fe-Mn, F e - M n - C r and F e - M n - C r - C alloys, the austenite-stable region for the F e - M n - C r - C system was determined [8]. The result of that work was a "modified Schaeffler diagram" for high manganese alloys. This diagram differed significantly from the standard Schaeffler diagram that was determined for high nickel steels containing only small amounts of manganese [8]. After determining the modified diagram, it was possible to pick a composition range where stable austenite, high manganese F e - M n - C r - C alloys should be obtained that could be further alloyed [8]. The studies suggested a composition range of Fe-( 12-14)Cr-(20-25)Mn-(0.1-0.25)C [8]. In this paper a base composition was chosen and the tensile properties were determined. That base composition was then alloyed for improved strength and for irradiation resistance by making additions of titanium, tungsten, vanadium, phosphorus and boron. With the exception of tungsten, these elements are known to be effective in nickel-based stainless steels [9-12]. Titanium and vanadium play roles in MC formation which can enhance strength. Tungsten was chosen as a substitute for molybdenum, which contributes strengthening to type 316 stainless steel but cannot be used in an FIRD steel. The effectiveness of boron and phosphorus has been attributed to an effect that these elements have on refining the M 2 3 C 6 precipitates that form in these
TABLE 1
types of steels. The objective was an FIRD alloy with strength and irradiation resistance equivalent to or better than that of type 316 stainless steel, which is considered a reference material for fusion reactor structural applications. Microstructural observations and tensile property determinations were made on several solutemodified alloys derived from the base composition, and the results from those studies will be presented here.
2. Experimentalprocedure On the basis of the previous work [8], a nominal Fe-20Mn-12Cr-0.25C steel was chosen as a base composition. Seven experimental alloys, including the base composition, were obtained in the form of 600 g vacuum-arc-melted button heats. Table 1 lists the alloys and their designations. By making elemental additions to the base composition (designated MnCrC), alloys were obtained with titanium (MnCrCTi), tungsten (MnCrCW), a combination of titanium and tungsten (MnCrCTiW) and combinations of these elements with vanadium, boron and phosphorus (MnCrCTiBP, M n C r C T i V B P and MnCrCTiWVBP). When additions of the various elements were made, nominal levels of titanium, tungsten, vanadium, boron and phosphorus of 0.1, 1, 0.1, 0.005 and 0.03 respectively were sought. Actual compositions are given in Table 1. Alloys were cast into ingots with a rectangular cross-section of 12.7 mm by 25.4 mm and a length of 152 mm. The ingots were hot rolled at 1050 °C to a thickness of approximately 6.4 mm. After homogenizing for 5 h at 1200 °C, the steel was cold rolled in five stages to sheet 0.76 mm thick. Between each stage the steel was annealed for 1 h at 1150°C. The sheet was finished in the
Fe-Cr-Mn alloys tested
Alloy designation MnCrC MnCrCTi MnCrCW MnCrCTiW MnCrCTiBP MnCrCTiVBP MnCrCTiWVBP aBalance is iron.
Composition (wt.% )a Cr
Mn
C
11.83 11.73 11.80 11.71 11.85 11.84 11.70
20.51 20.50 20.46 21.13 20.50 20.82 20.39
0.24 0.25 0.23 0.25 0.24 0.22 0.25
Ti
W
0.11
0.09 0.83 0.77
0.12 0.10 0.10 0.10
1.08
V
P
B
0.01 0.01 0.01 0.01 0.01 0.10 0.10
0.003 0.003 0.004 0.003 0.034 0.033 0.027
0.005 0.005 0.005
19 20% cold-worked condition. Tensile specimens were obtained from the cold-worked sheet, and specimens were subsequently heat treated to determine properties in the annealed condition. Tensile and transmission electron microscopy (TEM) specimens were inserted into a tube furnace already at the annealing temperature and annealed in flowing helium. They were then rapidly cooled by pulling them out of the furnace into a chamber that was cooled by flowing helium. The solution-annealed microstructures were examined by optical and electron microscopy. The TEM was performed using a Philips EM 430 (300 keV) and a JEOL 2000FX (200 keV). Some of the larger precipitate particles were identified in-foil using X-ray energy-dispersive spectroscopy (XEDS) capabilities on the 2000FX analytical electron microscope (AEM). Tensile tests on single specimens were made at room temperature, 200, 300, 400, 500 and 600 °C on specimens in the cold-worked condition and after two annealing treatments: 1 h at 1050°C and 2 h at 1150°C. Specimens had a reduced gauge section 20.3 mm long by 1.52 mm wide by 0.76 mm thick. All specimens were machined with gauge lengths parallel to the rolling direction. Tests were made in vacuum on a 120 kN capacity Instron universal testing machine at a cross-head speed of 8.5 x 10 -3 mm s-~, which produced a nominal strain rate of approximately 4.2 x 10-4 s- 1.
3. R e s u l t s and d i s c u s s i o n
Fig. 1. Opticalmicrostructuresof (a) type 316 stainlesssteel and (b) Fe-20Mn-12Cr-0.25C alloy annealed 1 h at 1050 °C.
3.1. O p t i c a l m i c r o s c o p y
Microstructures for the F e - 2 0 M n - 1 2 C r 0.25C base alloy (MnCrC) and type 316 stainless steel (316 SS) are shown in Fig. 1 after both were annealed for 1 h at 1050 °C. Overall microstructures for the two steels were similar: both appeared to contain scattered precipitates in the cross-section and had grain sizes estimated as ASTM number 4. When minor amounts of alloying elements were added to the base composition, the primary changes relative to the base alloy appeared to be in the amounts of precipitate present and the grain size (Table 2). Figure 2 shows the microstructures for the steels after annealing at 1050°C. A titanium addition (MnCrCTi) (Fig. 2(b)) refined the grain size and produced some
TABLE 2
ASTM grain size of steels tested
Alloy designation
MnCrC MnCrCTi MnCrCW MnCrCTiW MnCrCTiBP MnCrCTiVBP MnCrCTiWVBP 316 SS
Heat treatment 1 h at 1050 °C
2 h at 1150 °C
3 8 4 8 8 8 8 4
1 6 3 6 6 4 5
precipitation, whereas an addition of tungsten (MnCrCW) had only a small effect on the grain size and precipitation behavior relative to
20
Fig. 2. Optical microstructures of seven experimental heats of high manganese steels annealed 1 h at 1050 *C: (a) MnCrC; (b) MnCrCTi; (c) MnCrCW; (d) MnCrCTiW; (e) MnCrCTiBP; (f) MnCrCTiVBP; (g) MnCrCTiWVBP.
21
MnCrC (Fig. 2(c)). There was more precipitate present when titanium and tungsten were present in the same alloy (MnCrCTiW) (Fig. 2(d)). The amount of precipitate continued to increase as the alloying additions progressed to titanium, boron and phosphorus (MnCrCTiBP) (Fig. 2(e)),
to titanium, vanadium, boron and phosphorus (MnCrCTiVBP) (Fig. 2(f)) and to titanium, tungsten, vanadium, boron and phosphorus (MnCrCTiWVBP) (Fig. 2(g)). In addition to alloying effects, the sensitivity to changes in annealing temperature was also
Fig. 3. A comparison of the optical microstructures of (a), (b) MnCrCTiBP and (c), (d) MnCrCTiWVBP after annealing 1 h at 1050 *C ((a) and (c)) and 2 h at 1150 *C ((b) and (d)).
22
investigated. The primary difference between microstructures of steels annealed at 1050°C (Fig. 2) and those produced by annealing at 1150°C was that after the 1150°C anneal the major portion of the precipitate was found at the grain boundaries whereas at 1050°C precipitation within the matrix dominated. Typical microstructural differences are shown in Fig. 3 for t h e MnCrCTiBP and MnCrCTiWVBP alloys annealed at 1050 and 1150°C. In some cases at 1050 °C many of the precipitates appeared to be on grain boundaries, but on further examination these linear arrays of precipitates were found to be in the matrix. They apparently formed earlier at grain boundaries but the boundaries separated from the precipitates as grain growth continued (see Fig. 3(c)). This will be discussed further when the T E M results are presented. As shown in Table 2, grain size was considerably larger after the 115 0 °C anneal.
annealing at 1150 °C (Fig. 4). Electron diffraction confirmed the identification of M23C 6. When titanium was added to the base alloy (MnCrCTi), a relatively high density of finer (10-25 nm) and some coarser (90-600 nm) precipitates formed within the grains of the steel annealed at 1050 °C (Fig. 5). These precipitates
3.2. Electron microscopy Table 3 summarizes the precipitates observed in the manganese-stabilized stainless steels by TEM. No matrix precipitation was found in the MnCrC after 1 h at 1050°C or 2 h at 1150°C. Traces of fine (20-60 nm diameter) M23C 6 w e r e detected along the grain boundaries only after TABLE 3
f
-"
Fig. 4. Microstructure of the Fe-Mn-Cr-C (MnCrC) base alloy showing the M23C6 precipitates in grain boundary after annealing 2 h at 1150 °C.
Size and type of precipitates observed in manganese-stabilized stainless steels
Alloy designation
Precipitate description Annealed i h at 1050 °C Size (rim)a
MnCrC MnCrCTi
10-25(M) 90-600(M) No precipitate detected
MnCrCTiW
30-60 (M) 300-600 (M) 80-140 (M, GB) 300-1000 (M, GB) 15-60 (M) 50-100 (GB) 250-800 (GB) 15-60 (M) 200-1000 (M) 50-100 (GB) 250-800 (GB)
MnCrCTiWVBP
MnCrCTiWVBP
Type
No precipitate detected
MnCrCW
MnCrCTiBP
Annealed 2 h at 1150 °C
MC MC
MC MC MC MC MC MC Mz3C 6
MC MC MC
Size (am) a
20-60 (GB) 40-80 (GB) 90-600 (M)
Type M23C 6
MC MC
50-100 (M) 170-300 (GB)
M23C6 M23C6
400-600 (M)
MC, M23C6
80-140 (M, GB) 300-1000 (M, GB)
MC MC
300-1000 (M) 250-800 (GB) 50-100 (GB)
MC MC
300-1000 (M) 250-800 (GB)
MC, M23C6
M23C6
MC
M23C6
~Location of precipitate is given in parentheses after the size, where M represents matrix and GB represents grain boundary.
23
Fig, 5. Matrix precipitates (primarily MC) in MnCrCTi annealed 1 h at 1050 °(2.
were non-uniformly distributed spatially, with the smaller precipitates often appearing in rows or bands, which may indicate that the precipitates nucleated along boundaries that broke away and continued to migrate, either during recrystallization of the cold-worked structure or during grain growth. After annealing at 1150 °C, the MC precipitation (particles 40-80 nm in size) occurred at most of the grain boundaries (Fig. 6(a)), and only occasionally did large particles or clusters of smaller particles form within the grains. The MC at grain boundaries was confirmed by electron diffraction; n o M23C6 was detected. Tangled networks of dislocations surrounded the large MC particles or clusters found after annealing at 1150 °C (Fig. 6(b)). The base alloy with only tungsten added (MnCrCW) was essentially free of detectable precipitation after annealing at 1050 °C. Annealing at 1150°C did, however, produce a few M23C6 particles in the matrix. These M23C6 precipitates were generally not surrounded by the dislocation structure that surrounded the MC formed in the MnCrCTi. Electron diffraction indicated that only M23C6 precipitation (170-300 nm particles) formed on grain boundaries after annealing at 1150 °C; these were several times larger than the MC formed in the MnCrCTi (Fig. 7). Combined additions of titanium and tungsten to the base alloy (MnCrCTiW) produced a mixture of many fine (30-60 nm) and a few coarse
Fig. 6. (a) Fine MC precipitate in grain boundary and matrix precipitates in MnCrCTi and (b) the matrix precipitates at higher magnification after annealing 2 h at 1150 °C.
Fig. 7. Grain boundary M23C6 precipitates in MnCrCW annealed 2 h at 1150 °C.
24
Fig. 9. Microstructure of MnCrCTiBP annealed 1 h at 1050 °C showing matrix precipitate and dislocation structure.'
Fig. 8. Example of M23C 6 and MC matrix precipitation observed in the MnCrCTiW after annealing (a) 1 h at 1050 °C and (b) 2 h at 1150"(2.
(300-600 nm) MC precipitates in the matrix during annealing at 1050 °C (Fig. 8(a)). Annealing at 1150 °C produced only a coarse (400-600 nm) distribution of MC and M23C 6 particles. These were surrounded by tangles of dislocations (Fig. 8(b)). No grain boundary precipitation was detected in this alloy at either annealing temperature. Additions of phosphorus and boron in combination with titanium (MnCrCTiBP) produced a non-uniform distribution of some smaller (80-140 nm) and larger (0.3-1/~m) MC precipi-
tates, both in the matrix and at grain boundaries of the specimens annealed at both temperatures. The specimen annealed at 1050 °C had more dislocations and a higher concentration of annealing twins (Fig. 9) relative to the specimen annealed at 1150 °C. Many of the dislocations were within the planar twin boundaries or extended into the matrix from precipitate particles and appeared to have dissociated into partial dislocations with stacking faults in between. This dislocation configuration suggests that this structure developed during annealing or cooling. None of these features were observed after annealing at 1150°C. "Wavy" large angle grain boundaries were observed after both heat treatments; this morphology appears to be caused by faceting of the grain boundary (Fig. 10(a)). In many instances the facets appeared to be associated with precipitates along the grain boundaries (Fig. 10(b)). Vanadium added in combination with titanium, boron and phosphorus (MnCrCTiVBP) caused a large increase in the amount of fine (15-60 nm) intragranular MC precipitates that formed in the steel annealed at 1050 °C (compared to the steel without vanadium, boron and phosphorus). Many of these MC precipitates appeared in rows or
25
Fig. 11. Microstructure of MnCrCTiVBP annealed 1 h at 1050 °C showing fine MC and coarse M23C6 particles.
Fig. 10. Microstructure of MnCrCTiBP annealed 2 h at 1150 °C showing (a) "wavy" or faceted grain boundary and (b) higher magnification photomicrograph of the precipitates observed in association with wavy boundaries.
bands (Fig. 11 ). There was some grain boundary precipitation--mainly larger (250-800 nm) M23C 6 particles with possibly some smaller (50-100 nm) MC particles in between. After annealing at 1150°C, there was much less precipitation and most of the precipitation was on the grain boundaries; much more grain boundary precipitation was found at 1150 than at 1050 °C. Grain boundary faceting was observed after both annealing treatments. Finally, the addition of titanium, tungsten, vanadium, boron and phosphorus to the base composition (MnCrCTiWVBP) produced a relatively high density of intragranular MC precipitates at 1050 °C, many of them in rows or bands, as observed in the MnCrCTiVBP. As with the MnCrCTiVBP, a much lower density of matrix MC was found after annealing at 1150°C. Many of the particles found in the matrix were surrounded by dislocations. At 1150 °C the few MC precipitates that were observed were on the
grain boundaries, and again grain boundary faceting was evident (Fig. 12). Precipitates noted above were mainly identified in-foil by their characteristic electron diffraction patterns, but some precipitate particles in thin foil specimens annealed at 1150 °C were also analyzed by XEDS. Consistently, most of the coarse particles were chromium-rich M 2 3 C 6 and the finer particles in each sample were MC. Some coarse titanium-rich MC particles were identified in the matrix, but coarse particles at grain boundaries were M23C 6. With the exception of the MnCrC and the MnCrCW, more precipitate formed at 1050 than 1150 °C. This observation is probably due to the difference in carbide solubility at 1050 and 1150 °C. In normal 300 series stainless steels with about 0.05% C, the carbon is fully soluble at 950 and 1050 °C and above. However, in the M n - C r steels with 0.25% C, carbides can form even at 1150°C. The lower carbon solubility at 1050 than at 1150 °C should lead to a larger amount of carbide precipitate forming during the recrystallization and grain growth process at 1050 °C. Since precipitates are observed after the 1050 and 1150 °C anneals, these heat treatments are not true solution anneals.
26
Fig. 12. Grain boundary MC precipitation in the MnCrCTiWVBP annealed 2 h at 1150 °C.
Small TEM disks were used for the microstructural studies. They were cooled rapidly after annealing by pulling them from the furnace into an unheated chamber that was cooled by flowing helium. Such a rapid cooling would be expected to allow them to cool without precipitation occurring. However, the apparent anomalous observation of precipitate--especially grain boundary precipitate--appearing at 1150 and not at 1050 °C for the two alloys without titanium may mean that no precipitate was present when annealed at 1050 and 1150 °C, but precipitation occurred in the steel annealed at 1150 °C because of the slightly longer time that is necessary for a specimen to cool to room temperature from 1150 than from 1050°C. If these precipitates were present at the annealing temperature, an alternative explanation is that the apparent anomaly is due to the dynamics of grain boundary precipitation. That is, after 1 h at 1050 °C, grain boundary motion from recrystallization and/or grain growth processes is continuing, whereas an equilibrium grain size is probably approached in the 2 h at 1150°C. For a system only slightly supersaturated, a precipitate can more easily nucleate on a stationary boundary than a moving one. Matrix precipitates that form in rows or bands after the 1050 °C anneal in several of the other alloys suggest that the precipitation at 1050 °C is intimately connected with the recrystallization process.
3.3. Tensile behavior Figures 13 and 14 show a comparison of the tensile properties of type 316 SS with the Fe-20Mn-12Cr-0.25C base alloy after annealing for 1 h at 1050°C and after being cold worked 20% (common conditions in which such steels are used). The yield stress (YS) of the manganese-stabilized steel in both conditions was equivalent to that of 316 SS (Fig. 13(a)). Since manganese additions cause an increased workhardening rate, the MnCrC steel reached a higher ultimate tensile strength (UTS) when tested in both thermomechanical conditions (Fig. 13(b)). Despite this higher work-hardening capability, the high manganese steel still had equivalent or better ductility (as measured by total elongation) in both the annealed and cold-worked conditions (Fig. 14). These tensile results indicated that an austenitic base alloy with substantial strength and ductility could be obtained by using the information developed from the modified Schaeffier diagram [8]. The next objective was to improve the strength of the new manganese-stabilized alloys. This was attempted by adding titanium, tungsten, vanadium, boron and phosphorus in the quantities given in Table 1 to the F e - 2 0 M n - 1 2 C r 0.25C base composition. Table 4 gives the room temperature tensile properties for the modified alloys, the base alloy and 316 SS as a reference. The manganese-
27 900
'
800 -
1
'
I
I ~ I ~ C O L D
'
I
WORKED
--~
Alloy
700
~ 600 500
MnCrC
• •
-_
316 SS
20O --i~"--~.,.,,l!,~__¢~::::=~ S0 L UT I0 N ANNEALED _
i
I
I
L
200
(a)
i
I
400
600
TEMPERATURE (*C) I100
B
i
I
I
[ •
I000 900
~
800
~
.
i
MnCrC
• 5t6
-r
z
C O L D ~ WORKED ~.
~
-
SS
_ ~l
L-
700 z uJ I./ bJ 6 0 0 r--e-~-
500
SOLUTION ~ ANNEALED ~N
•
-
~l
|
_
.-------.,,,
4OO I
I
I 2o0
0
i
i 40o
i
-'~'600
TEMPERATURE (°C)
(b)
Fig. 13. (a) 0.2% yield stress and (b) ultimate tensile strength plotted against test temperature for the F e - M n - C r - C alloy and type 316 stainless steel in the 20% cold-worked condition and annealed 1 h at 1050 °C.
7o
i
I
~a------~. k
~
ANNEALED
0
• MnCrC
-J 50 --
• 316 SS
Ld J
\
N I / I
i
I
I
I -q
~ 20 - - i ~
~,'COLD WORKED
io o
Elongation (%)
YS
Uniform
Total
UTS
~.,'
300
I00
Strength (MPa)
designation -
400 w
TABLE 4 Room temperature tensile properties of manganese-stabilized steels and type 3 1 6 stainless steel
I
4 t
I
20O 400 TEMPERATURE (*C)
I 600
Fig. 14. Total elongation plotted against test temperature for the F e - M n - C r - C (MnCrC) alloy and type 316 stainless steel in the 20% cold-worked condition and annealed 1 h at 1050 °C.
stabilized steels were tested in two annealed conditions (1 h at 1050 °C and 2 h at 1150°C) and in the 20% cold-worked condition. The 316 SS was tested after the 1050 °C solution anneal and after cold working 20%.
Annealed 1 h at 1050 °C MnCrC 220 MnCrCTi 279 MnCrCW 267 MnCrCTiW 302 MnCrCTiBP 288 MnCrCTiVBP 275 MnCrCTiWVBP 304 316 SS 236
798 927 803 918 935 935 915 586
55.4 49.7 57.1 53.8 52.2 51.0 54.9 54.3
56.6 53.0 59.9 56.9 55.6 53.9 57.5 58.2
Annealed 2 h at 1150 °C MnCrC 233 MnCrCTi 258 MnCrCW 247 MnCrCTiW 258 MnCrCTiBP 271 MnCrCTiVBP 221 MnCrCTiWVBP 264
766 891 761 882 891 859 869
53.4 53.5 55.4 54.5 52.8 49.6 59.9
55.1 56.4 57.0 57.2 54.2 50.4 61.7
1086 1160 1057 1168 1158 1126 1114 807
14.1 10.7 17.6 6.6 10.4 11.4 11.3 11.5
16.0 13.0 20.0 9.5 12.1 13.1 13.6 17.4
20% coM worked MnCrC MnCrCTi MnCrCW MnCrCTiW MnCrCTiBP MnCrCTiVBP MnCrCTiWVBP 316 SS
815 954 784 980 946 862 915 739
The room temperature results indicated that the strength of a given high manganese steel annealed for 1 h at 1050 °C generally exceeded that same steel annealed for 2 h at 1150 °C. This behavior probably reflects both more and finer precipitation and a finer grain size for specimens annealed at the lower temperature. With one exception, the YS and UTS of the manganesestabilized stainless steels annealed at 1050°C exceeded those of the 316 SS annealed at this temperature. The exception was that the YS for the MnCrC steel annealed at 1050 °C was slightly lower than the YS of 316 SS. (Note also that only two of the YS values for the manganese-stabilized steels annealed at 1150 °C were less than that of 316 SS annealed at 1050 °C.) The YS and UTS of all of the manganese-stabilized steels in the 20% cold-worked condition exceeded those for 316 SS in that condition. In both annealed conditions the uniform and total elongations of the manganese-stabilized steels at room temperature were similar to those for 316 SS (Table 4). Similar ductility was also observed for most of the alloys in the coldworked condition. The only exception was the MnCrCTiW alloy, which had the lowest uniform
28
and total elongations, although the observed values still indicate adequate ductility. The addition of tungsten (MnCrCW) to the base composition had the least effect on the strength of the annealed steels. An addition of titanium by itself (MnCrCTi) had only a modest strengthening effect. A combination of titanium, boron and phosphorus (MnCrCTiBP) and these three elements plus vanadium (MnCrCTiVBP) gave rise to strengths that were similar to those of MnCrCTi. However, the combination of titanium and tungsten (MnCrCTiW) produced a substantial strength increase over the base' composition. Adding vanadium, boron and phosphorus to the combination of tungsten and titanium did not significantly change the room temperature strength in the annealed condition from that found in MnCrCTiW. In the cold-worked condition tested at room temperature, trends in the relative strength observed for the various combination of elements were similar to those observed in the annealed steels, with the exception that the MnCrCTiW had higher YS and UTS values than the MnCrCTiWVBE However, this greater strength
was also accompanied by lower uniform and total elongation values. Tensile tests were conducted over the range from room temperature to 600 °C. In Figs. 15-20 the tensile properties for the five strongest manganese-stabilized steels (MnCrC and MnCrCW are not shown) are compared with type 316 SS. 60
os4~. ~
48
w ,,4
42
~ ,-0,~ •-0-..-0-
36
30 0
MnCrCTiW X MnCrCTiPB ~) MnCrCTiVPB MnCrCTiWVPB 316 SS I I r I 100 200 300 400 TEMPERATURE(=C)
~ . ~
\ I 500
600
Fig. 16. Total elongation plotted against test temperature for five heats of manganese-stabilized stainless steels and type 316 stainless steel as a function of test temperature. All steels were annealed 1 h at 1050 °C.
270 320
~ 0~. ~ ~..~ ~
280
~ MnCrCTI ~ MnCrCTiW -0-- MnCrCTiPB " 0 - MnCrCTIVPB --o-- MnCrCTiWVPB
~ ~ ~
240
~
MnCrCTi MnCrCTIW
t
240 A
MnCrCTiWVPB.~
~ 210 ~ 180
~ X
200
9 uJ
160
9
~. 150 120
120
9O I
80
I
I
1O0
I
I
200 300 400 TEMPERATURE(°C)
(a)
500
600
1O0
200
(a)
300 400 TEMPERATURE(°C)
500
600
900 1ooo
~ ~--~=~ ~
900
800
2
-.e- MnCrCTi ~ MnCrCTiW ,-o- MnCrCTIP6
~
\~-%
'
850
--Z..OrOT,' MnCrCTiW MnCrCTiPB M~CrrC;:V:?pB
coo
750
~,oT,vP.
700
uJ 650 700
~
S
S
Z
600
uJ 550 -~ h
500
450
~ 500
400 400 0 (b)
I 100
I I I 200 300 400 TEMPERATURE(=C)
I 500
600
Fig. 15. (a) 0.2% yield stress and (b) ultimate tensile strength plotted against test temperature for five heats of manganesestabilized stainless steels and type 316 stainless steel. All steels were annealed 1 h at 1050 °C.
(b)
i
1O0
1
i
i
200 300 400 TEMPERATURE(°C)
i
500
600
Fig. 17. (a) 0.2% yield stress and (b) ultimate tensile strength plotted against test temperature for five heats of manganesestabilized stainless steels and type 316 stainless steel. The manganese steels were annealed 2 h at 1150 °C and the type 316 stainless steel was annealed 1 h at 1050 °C.
29 18
64
z _0
,
6O
16
se
14
i ¸
k ~ _ ~
i
i
i
•
MnCrCTi ~ MnCrCTIW •--o--- MnCrCTiPB ~ MnCrCTIVPB ---o--- MnCrCTiWVPB
~ ~
52
q
z
10
w ~
44
uJ j
8
~
4o
~
6
36
4
32
0
100
200 300 400 TEMPERATURE (°C)
500
600
1000
900
~ ~ ~ ~ .-0.,-
MnCrCTi MnCrCTiW MnCrCTiPB MnCrCTiVPB MnCrCTiWVPB
- - C - - 3 1 6 SS
800
00
~0 700 uJ >6O0
500
I 100
0 (a)
1200 11oo
I I I 200 300 400 TEMPERATURE (*C)
~
•
~
~ -o~ ~
~ . , . ~ ~ . ' ~ ~
(~'r" 1000
°
I 500
600
MnCrCTi MnCrCTiW ~CrCTipB MnCrCTiVPB MnCrCTiWVPB
900 ¢n Z
800
i
700 600
I
0
(b)
100
I 200
I I 300 400 TEMPERATUREeC)
0
I 100
I 200
I 300
I 400
I 500
600
TEMPERATURE (°C)
Fig. 18. Total elongation plotted against test temperature for five heats of manganese-stabilized stainless steels and type 316 stainless steel. The manganese steels were annealed 2 h at 1150 °C and the type 316 stainless steel was annealed 1 h at 1050 °C.
.~ " ~ , . ~,,,~'~= . t : ~ ~ ~'~---~ " ",~'111~
2
I 500
600
Fig. 19. (a) 0.2% yield stress and (b) ultimate tensile strength plotted against test temperature for five heats of manganesestabilized stainless steels and type 316 stainless steel. All steels were in the 20% cold-worked condition.
Results clearly indicated a superiority in strength for the manganese-stabilized steels relative to 316 SS after an anneal for 1 h at 1050°C (Figs. 15 and 16), 2 h at 1150 °C (Figs. 17 and 18) and
Fig. 20. Total elongation plotted against test temperature for five heats of manganese-stabilized stainless steels and type 316 stainless steel. All steels were in the 20% cold-worked condition.
after 20% cold work (Figs. 19 and 20). Despite their greater strength, the ductility of the manganese-stabilized steels was generally equivalent to or better than that for 316 SS in the annealed condition (Figs. 16, 18 and 20). These observations on ductility are important because they mean that these new steels represent a significant gain in strength that does not come at the expense of ductility. For the steels annealed for 1 h at 1 0 5 0 ° C the MnCrCTiW and the M n C r C T i W V B P had the highest YS of the modified steels (Fig. 15(a)), with the MnCrCTiWVBP being the strongest at the highest test temperatures. Similar UTS behavior was observed (Fig. 15(b)), with the MnCrCTiWVBP being essentially the strongest alloy over most of the temperature range. Ductilities of all the manganese-stabilized steels annealed at 1050°C were excellent over the entire temperature range and were generally equivalent to or better than those for type 316 SS (Fig. 16). Tensile data for the steels annealed for 2 h at 1150°C (Figs. 17 and 18) displayed the same general behavior as those annealed for 1 h at 1050°C. The MnCrCTiWVBP again had the best strength properties at the highest test temperatures and also had excellent ductility relative to the other steels and type 316 SS. In the cold-worked condition the YS of the MnCrCTiW was highest over most of the temperature range (Fig. 19(a)). The microstructure of this alloy developed stacking faults and dissociated dislocations when aged, suggesting that it has the lowest stacking fault energy of the modified alloys [12], which may contribute to its
30 strength. There was much less difference in the UTS behavior of the five manganese-stabilized steels over this temperature range (Fig. 19(b)). Both YS and UTS of all the modified manganesestabilized steels were substantially greater than found in 20% cold-worked 316 SS. The ductility (Fig. 20) of the MnCrCTiW was lowest and that of the MnCrCTiWVBP was highest for much of the temperature range, although the MnCrCTiWVBP had the lowest ductility at 200 °C. All of the cold-worked F e - M n - C r steels had less ductility than cold-worked 316 SS from room temperature to 200 °C, but then most or all had better ductility from 300 to 600 °C. Fractures for all thermomechanical treatments were ductile and accompanied by necking. The minimum in ductility around 2 0 0 - 4 0 0 °C for the cold-worked, high manganese steels is also observed for cold-worked nickel-containing stainless steels, as seen in Fig. 20 for type 316 stainless steel. Detailed study of this behavior for the fusion prime candidate alloy (PCA), a 16Ni-14Cr-2.5Mo-0.08C austenitic alloy that contains small amounts of titanium, niobium, vanadium, nitrogen and phosphorus, indicated that a cup-cone failure occurred for the low and high temperature tests [13]. A ductile-shear-type fracture with the fracture surface at about 35 ° to the tensile axis occurred at temperatures near the minimum. Scanning electron microscopy studies indicated that this minimum is related to a change from void coalescence at the low and elevated test temperatures to ductile tearing at temperatures of the minimum [13]. These results indicate that it is possible to develop an FIRD steel free of nickel, molybdenum, nitrogen and copper, with tensile properties that are better than those of type 316 stainless steel. In fact, the room temperature YS of the 20% cold-worked MnCrCTiWVBP exceeded that of the PCA [11]. However, the rate at which the YS of the MnCrCTiWVBP decreases with temperature is greater than for the PCA, and above about 450 °C the PCA becomes stronger. The PCA has been under development for several years to optimize its properties and microstructural behavior for strength and irradiation resistance. Further optimization of the properties of the manganese steels should also be possible by adjusting the solute additions and thermomechanlcal treatment to affect the solid solution matrix and the precipitate microstructure.
The tensile results generally indicate the importance of minor alloying additions, e.g. titanium for strength, and synergistic effects make combinations of additives, e.g. titanium and tungsten, more important than single-element additions (Table 4 and Figs. 15-20). Steels containing titanium were generally among the strongest; these were basically the same steels that contained fine MC matrix precipitates observed by T E M in steels annealed at 1050°C in this work or in cold-worked steels aged at 800 °C in previous work [12]. The steel with the combination of vanadium and titanium but without any tungsten (MnCrCTiVBP) may also be of interest. Other work shows the MnCrCTiVBP to be very resistant to recrystallization and the formation of intermetallic phases, e.g. a and Laves, when aged at 800 °C in the cold-worked condition [12]. This steel contained little precipitation after annealing, even though vanadium would be expected to react in conjunction with titanium to form MC. The MnCrCTiVBP was generally the weakest of the steels in the annealed condition, but in the cold-worked condition it was similar to the others at higher temperatures. Tungsten additions to this combination of elements definitely improved the strength in both the annealed and the coldworked conditions. This observation emphasizes the fact that synergistic effects of tungsten and titanium appear to be important. Although the experimental steels discussed in this paper have good strength, they must also be weldable and be compatible with the coolant used in the blanket structure in order to be used for building fusion reactors. Corrosion and weldability studies are in progress. The steels are also being irradiated to determine their resistance to irradiation-induced swelling.
4. Summary and conclusions A series of experimental, high manganese austenitic stainless steels were produced. An F e - 2 0 M n - 1 2 C r - 0 . 2 5 C steel was used as a base alloy composition which was then alloyed for improved strength by adding small amounts of titanium, tungsten, vanadium, boron and phosphorus. Austenite stabilization was achieved with manganese and carbon alone, without the use of high nickel and/or high nitrogen concentrations, as is the case in most high manganese austenitic steels. In tensile tests over the range from room temperature to 600 °C the manganese-stabilized
31
steels had 0.2% yield stress and ultimate tensile strength values that were superior to those of type 316 stainless steel, a commercially available nickel-stabilized austenitic stainless steel. The manganese-stabilized steels were superior in both the solution-annealed and cold-worked conditions. The ductility of the manganese steels was as good as or better than that of the 316 stainless steel for most of the test conditions. Acknowledgments We wish to thank the following people who helped in the completion of this work: N. H. Rouse carried out the tensile tests and prepared the electron microscopy specimens; E. A. Kenik and M. L. Santella reviewed the manuscript; Frances Scarboro prepared the manuscript. This research is sponsored by the Office of Fusion Energy, U.S. Department of Energy, under Contract No. DE-AC05-840R21400 with the Martin Marietta Energy Systems, Inc. References 1 R. L. Klueh and E. E. Bloom, in F. A. Garner, D. S. Gelles and E W. Wiffen (eds.), Optimizing Materials for
2 3 4 5 6 7
8 9 10 11
Nuclear Applications, The Metallurgical Society, Warrendale, PA, 1985, pp. 73-85. J. H. G. Monypenny, Stainless Iron and Steel Vol. 2, Chapman and Hall, London, 1954, p. 261. E. Ruedl, D. Rickerby and T. Sasaki, Fusion Technology, Vol. 2, Pergamon, Oxford, 1984, p. 1029. E. Ruedl and T. Sasaki, J. Nucl. Mater., 122-123 (1984) 794-798. H.R. Brager, F. A. Garner, D. S. Gelles and M. L. Hamilton, J. Nucl. Mater., 133-134 (1985) 907-911. A . H . Bott, F. B. Picketing and G. J. Butterworth, J. NucL Mater., 141-143(1986) 1088-1096. M. Tamura, H. Hayakawa, M. Tanimura, A. Hishinuma and T. Kondo, J. NucL Mater., 141-143 (1986) 1067-1073. R. L. Klueh, P. J. Maziasz and E. H. Lee, Mater. Sci. Eng., A102 (1988) 115-124. P. J. Maziasz and T. K. Roche, J. NucL Mater., 103-104 (1981) 797-802. P. J. Maziasz et aL, J. Nucl. Mater., 108-109 (1982) 296-298. P.J. Maziasz, in B. L. Bramfitt, R. L. Benn, C. R. Brinkman and G. E Vander Voort (eds.), Optimization of Processing, Properties, and Service Performance Through Microstructural Control, American Society for Testing
and Materials, Philadelphia, PA, 1988, pp. 116-161. 12 P.J. Maziasz and R. L. Klueh, in R. L. Klueh, D. S. Gelles, M. Okada and N. H. Packan (eds.), Reduced Activation
Materials for Fusion Reactor Applications, ASTM Spec. Tech. Publ., 1047 (1990) 56-79. 13 D. N. Braski and P. J. Maziasz, J. Nucl. Mater., 122-123 (1984) 338-342.
17
Tensile and Microstructural Behavior of Solute-modified Manganese-stabilized Austenitic Steels R. L. KLUEH and P. J. MAZIASZ Metals and Ceramics Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6376 (U.S.A.) (Received September 5, 1989; in revised form December 4, 1989)
Abstract
A program is under way to develop a nickelfree austenitic stainless steel for fusion reactor applications. Previous work showed that a stable austenite alloy is possible in a compositional regime of Fe-(20-25)Mn-(12-14)Cr-(0.1-0.25)C. As a base composition for further alloying for strength and irradiation resistance, an Fe-20Mn12Cr-0.25C composition was chosen. Tensile properties for this base composition were comparable to those of type 316 stainless steel. To improve strength and irradiation resistance, closely controlled quantities of titanium, tungsten, vanadium, phosphorus and boron were added to the base alloy. Various combinations of these alloying elements improved the tensile properties of the new steels relative to type 316 stainless steel in both the annealed and 20% cold-worked conditions. These solute combinations also produced MC and M23 C6 precipitation in some alloys during annealing at 1050 and 1150 °C.
I. Introduction
Elements that make up the alloys proposed for the first wall and blanket structure of a magnetic fusion reactor become activated by the high energy neutrons created by the fusion reaction in the plasma. These radioactive reactor components must be properly disposed of after their service lifetime. The complexity of the waste disposal procedure depends on the time required for the induced radioactivity to decay to levels that are no longer hazardous to people and the environment. The more rapidly the radioactivity decays, the simpler the disposal task becomes. Part 61 of Title 10 of the U.S. Code of Federal Regulations (10CFR61) was prepared by the Nuclear Regulatory Commission for radioactive 0921-5093/90/$3.50
fission product disposal, Using 10CFR61, it is possible to define compositional limits for alloys that can be disposed of b y relatively simple shallow land burial techniques, as opposed to the much more difficult and expensive deep geologic burial. The objective of the present work is to develop new alloys that meet guidelines for shallow land burial. Some common alloying elements that decay over a long period of time after neutron activation are nickel, molybdenum, nitrogen, copper and niobium. An alloy development program is in progress [1] to develop fast induced radioactivity decay (FIRD) austenitic stainless steels to replace the high nickel stainless steels that are among the present candidate alloys for fusion applications. To produce an FIRD austenitic steel, manganese was proposed as a replacement for nickel, although it was recognized that this may prove difficult, because manganese is not as strong an austenite stabilizer as nickel and because the metallurgical effects of nickel and manganese in the alloy may not otherwise be equivalent [1]. The first attempts to use manganese as a replacement for nickel in stainless steels were made in the 1930s [2]. Since manganese and nickel behave differently in austenitic steels, commercially available high manganese stainless steels contain some nickel and use high concentrations of nitrogen to produce austenite stability. Other groups around the world are also working to develop high manganese, reduced activation austenitic steels [3-7]. However, in much of that work nitrogen is used for austenite stabilization [3, 4, 6] and in other instances small amounts of nickel are allowed [3, 4]. In keeping with the effort to meet the criteria established by the 10CFR61 guidelines, neither nitrogen nor nickel was used for austenite stabilization in the present work. © Elsevier Sequoia/Printed in The Netherlands
18
In the strategy proposed for developing FIRD steels [1], the first task was a determination of F e - M n - C r - C compositions that would produce a fully austenitic microstructure [1]. Such a base composition could then be modified with minor element additions to obtain strength and irradiation resistance at least equal to that found in current candidate fusion reactor structural materials. By investigating the microstructures of a series of Fe-Mn, F e - M n - C r and F e - M n - C r - C alloys, the austenite-stable region for the F e - M n - C r - C system was determined [8]. The result of that work was a "modified Schaeffler diagram" for high manganese alloys. This diagram differed significantly from the standard Schaeffler diagram that was determined for high nickel steels containing only small amounts of manganese [8]. After determining the modified diagram, it was possible to pick a composition range where stable austenite, high manganese F e - M n - C r - C alloys should be obtained that could be further alloyed [8]. The studies suggested a composition range of Fe-( 12-14)Cr-(20-25)Mn-(0.1-0.25)C [8]. In this paper a base composition was chosen and the tensile properties were determined. That base composition was then alloyed for improved strength and for irradiation resistance by making additions of titanium, tungsten, vanadium, phosphorus and boron. With the exception of tungsten, these elements are known to be effective in nickel-based stainless steels [9-12]. Titanium and vanadium play roles in MC formation which can enhance strength. Tungsten was chosen as a substitute for molybdenum, which contributes strengthening to type 316 stainless steel but cannot be used in an FIRD steel. The effectiveness of boron and phosphorus has been attributed to an effect that these elements have on refining the M 2 3 C 6 precipitates that form in these
TABLE 1
types of steels. The objective was an FIRD alloy with strength and irradiation resistance equivalent to or better than that of type 316 stainless steel, which is considered a reference material for fusion reactor structural applications. Microstructural observations and tensile property determinations were made on several solutemodified alloys derived from the base composition, and the results from those studies will be presented here.
2. Experimentalprocedure On the basis of the previous work [8], a nominal Fe-20Mn-12Cr-0.25C steel was chosen as a base composition. Seven experimental alloys, including the base composition, were obtained in the form of 600 g vacuum-arc-melted button heats. Table 1 lists the alloys and their designations. By making elemental additions to the base composition (designated MnCrC), alloys were obtained with titanium (MnCrCTi), tungsten (MnCrCW), a combination of titanium and tungsten (MnCrCTiW) and combinations of these elements with vanadium, boron and phosphorus (MnCrCTiBP, M n C r C T i V B P and MnCrCTiWVBP). When additions of the various elements were made, nominal levels of titanium, tungsten, vanadium, boron and phosphorus of 0.1, 1, 0.1, 0.005 and 0.03 respectively were sought. Actual compositions are given in Table 1. Alloys were cast into ingots with a rectangular cross-section of 12.7 mm by 25.4 mm and a length of 152 mm. The ingots were hot rolled at 1050 °C to a thickness of approximately 6.4 mm. After homogenizing for 5 h at 1200 °C, the steel was cold rolled in five stages to sheet 0.76 mm thick. Between each stage the steel was annealed for 1 h at 1150°C. The sheet was finished in the
Fe-Cr-Mn alloys tested
Alloy designation MnCrC MnCrCTi MnCrCW MnCrCTiW MnCrCTiBP MnCrCTiVBP MnCrCTiWVBP aBalance is iron.
Composition (wt.% )a Cr
Mn
C
11.83 11.73 11.80 11.71 11.85 11.84 11.70
20.51 20.50 20.46 21.13 20.50 20.82 20.39
0.24 0.25 0.23 0.25 0.24 0.22 0.25
Ti
W
0.11
0.09 0.83 0.77
0.12 0.10 0.10 0.10
1.08
V
P
B
0.01 0.01 0.01 0.01 0.01 0.10 0.10
0.003 0.003 0.004 0.003 0.034 0.033 0.027
0.005 0.005 0.005
19 20% cold-worked condition. Tensile specimens were obtained from the cold-worked sheet, and specimens were subsequently heat treated to determine properties in the annealed condition. Tensile and transmission electron microscopy (TEM) specimens were inserted into a tube furnace already at the annealing temperature and annealed in flowing helium. They were then rapidly cooled by pulling them out of the furnace into a chamber that was cooled by flowing helium. The solution-annealed microstructures were examined by optical and electron microscopy. The TEM was performed using a Philips EM 430 (300 keV) and a JEOL 2000FX (200 keV). Some of the larger precipitate particles were identified in-foil using X-ray energy-dispersive spectroscopy (XEDS) capabilities on the 2000FX analytical electron microscope (AEM). Tensile tests on single specimens were made at room temperature, 200, 300, 400, 500 and 600 °C on specimens in the cold-worked condition and after two annealing treatments: 1 h at 1050°C and 2 h at 1150°C. Specimens had a reduced gauge section 20.3 mm long by 1.52 mm wide by 0.76 mm thick. All specimens were machined with gauge lengths parallel to the rolling direction. Tests were made in vacuum on a 120 kN capacity Instron universal testing machine at a cross-head speed of 8.5 x 10 -3 mm s-~, which produced a nominal strain rate of approximately 4.2 x 10-4 s- 1.
3. R e s u l t s and d i s c u s s i o n
Fig. 1. Opticalmicrostructuresof (a) type 316 stainlesssteel and (b) Fe-20Mn-12Cr-0.25C alloy annealed 1 h at 1050 °C.
3.1. O p t i c a l m i c r o s c o p y
Microstructures for the F e - 2 0 M n - 1 2 C r 0.25C base alloy (MnCrC) and type 316 stainless steel (316 SS) are shown in Fig. 1 after both were annealed for 1 h at 1050 °C. Overall microstructures for the two steels were similar: both appeared to contain scattered precipitates in the cross-section and had grain sizes estimated as ASTM number 4. When minor amounts of alloying elements were added to the base composition, the primary changes relative to the base alloy appeared to be in the amounts of precipitate present and the grain size (Table 2). Figure 2 shows the microstructures for the steels after annealing at 1050°C. A titanium addition (MnCrCTi) (Fig. 2(b)) refined the grain size and produced some
TABLE 2
ASTM grain size of steels tested
Alloy designation
MnCrC MnCrCTi MnCrCW MnCrCTiW MnCrCTiBP MnCrCTiVBP MnCrCTiWVBP 316 SS
Heat treatment 1 h at 1050 °C
2 h at 1150 °C
3 8 4 8 8 8 8 4
1 6 3 6 6 4 5
precipitation, whereas an addition of tungsten (MnCrCW) had only a small effect on the grain size and precipitation behavior relative to
20
Fig. 2. Optical microstructures of seven experimental heats of high manganese steels annealed 1 h at 1050 *C: (a) MnCrC; (b) MnCrCTi; (c) MnCrCW; (d) MnCrCTiW; (e) MnCrCTiBP; (f) MnCrCTiVBP; (g) MnCrCTiWVBP.
21
MnCrC (Fig. 2(c)). There was more precipitate present when titanium and tungsten were present in the same alloy (MnCrCTiW) (Fig. 2(d)). The amount of precipitate continued to increase as the alloying additions progressed to titanium, boron and phosphorus (MnCrCTiBP) (Fig. 2(e)),
to titanium, vanadium, boron and phosphorus (MnCrCTiVBP) (Fig. 2(f)) and to titanium, tungsten, vanadium, boron and phosphorus (MnCrCTiWVBP) (Fig. 2(g)). In addition to alloying effects, the sensitivity to changes in annealing temperature was also
Fig. 3. A comparison of the optical microstructures of (a), (b) MnCrCTiBP and (c), (d) MnCrCTiWVBP after annealing 1 h at 1050 *C ((a) and (c)) and 2 h at 1150 *C ((b) and (d)).
22
investigated. The primary difference between microstructures of steels annealed at 1050°C (Fig. 2) and those produced by annealing at 1150°C was that after the 1150°C anneal the major portion of the precipitate was found at the grain boundaries whereas at 1050°C precipitation within the matrix dominated. Typical microstructural differences are shown in Fig. 3 for t h e MnCrCTiBP and MnCrCTiWVBP alloys annealed at 1050 and 1150°C. In some cases at 1050 °C many of the precipitates appeared to be on grain boundaries, but on further examination these linear arrays of precipitates were found to be in the matrix. They apparently formed earlier at grain boundaries but the boundaries separated from the precipitates as grain growth continued (see Fig. 3(c)). This will be discussed further when the T E M results are presented. As shown in Table 2, grain size was considerably larger after the 115 0 °C anneal.
annealing at 1150 °C (Fig. 4). Electron diffraction confirmed the identification of M23C 6. When titanium was added to the base alloy (MnCrCTi), a relatively high density of finer (10-25 nm) and some coarser (90-600 nm) precipitates formed within the grains of the steel annealed at 1050 °C (Fig. 5). These precipitates
3.2. Electron microscopy Table 3 summarizes the precipitates observed in the manganese-stabilized stainless steels by TEM. No matrix precipitation was found in the MnCrC after 1 h at 1050°C or 2 h at 1150°C. Traces of fine (20-60 nm diameter) M23C 6 w e r e detected along the grain boundaries only after TABLE 3
f
-"
Fig. 4. Microstructure of the Fe-Mn-Cr-C (MnCrC) base alloy showing the M23C6 precipitates in grain boundary after annealing 2 h at 1150 °C.
Size and type of precipitates observed in manganese-stabilized stainless steels
Alloy designation
Precipitate description Annealed i h at 1050 °C Size (rim)a
MnCrC MnCrCTi
10-25(M) 90-600(M) No precipitate detected
MnCrCTiW
30-60 (M) 300-600 (M) 80-140 (M, GB) 300-1000 (M, GB) 15-60 (M) 50-100 (GB) 250-800 (GB) 15-60 (M) 200-1000 (M) 50-100 (GB) 250-800 (GB)
MnCrCTiWVBP
MnCrCTiWVBP
Type
No precipitate detected
MnCrCW
MnCrCTiBP
Annealed 2 h at 1150 °C
MC MC
MC MC MC MC MC MC Mz3C 6
MC MC MC
Size (am) a
20-60 (GB) 40-80 (GB) 90-600 (M)
Type M23C 6
MC MC
50-100 (M) 170-300 (GB)
M23C6 M23C6
400-600 (M)
MC, M23C6
80-140 (M, GB) 300-1000 (M, GB)
MC MC
300-1000 (M) 250-800 (GB) 50-100 (GB)
MC MC
300-1000 (M) 250-800 (GB)
MC, M23C6
M23C6
MC
M23C6
~Location of precipitate is given in parentheses after the size, where M represents matrix and GB represents grain boundary.
23
Fig, 5. Matrix precipitates (primarily MC) in MnCrCTi annealed 1 h at 1050 °(2.
were non-uniformly distributed spatially, with the smaller precipitates often appearing in rows or bands, which may indicate that the precipitates nucleated along boundaries that broke away and continued to migrate, either during recrystallization of the cold-worked structure or during grain growth. After annealing at 1150 °C, the MC precipitation (particles 40-80 nm in size) occurred at most of the grain boundaries (Fig. 6(a)), and only occasionally did large particles or clusters of smaller particles form within the grains. The MC at grain boundaries was confirmed by electron diffraction; n o M23C6 was detected. Tangled networks of dislocations surrounded the large MC particles or clusters found after annealing at 1150 °C (Fig. 6(b)). The base alloy with only tungsten added (MnCrCW) was essentially free of detectable precipitation after annealing at 1050 °C. Annealing at 1150°C did, however, produce a few M23C6 particles in the matrix. These M23C6 precipitates were generally not surrounded by the dislocation structure that surrounded the MC formed in the MnCrCTi. Electron diffraction indicated that only M23C6 precipitation (170-300 nm particles) formed on grain boundaries after annealing at 1150 °C; these were several times larger than the MC formed in the MnCrCTi (Fig. 7). Combined additions of titanium and tungsten to the base alloy (MnCrCTiW) produced a mixture of many fine (30-60 nm) and a few coarse
Fig. 6. (a) Fine MC precipitate in grain boundary and matrix precipitates in MnCrCTi and (b) the matrix precipitates at higher magnification after annealing 2 h at 1150 °C.
Fig. 7. Grain boundary M23C6 precipitates in MnCrCW annealed 2 h at 1150 °C.
24
Fig. 9. Microstructure of MnCrCTiBP annealed 1 h at 1050 °C showing matrix precipitate and dislocation structure.'
Fig. 8. Example of M23C 6 and MC matrix precipitation observed in the MnCrCTiW after annealing (a) 1 h at 1050 °C and (b) 2 h at 1150"(2.
(300-600 nm) MC precipitates in the matrix during annealing at 1050 °C (Fig. 8(a)). Annealing at 1150 °C produced only a coarse (400-600 nm) distribution of MC and M23C 6 particles. These were surrounded by tangles of dislocations (Fig. 8(b)). No grain boundary precipitation was detected in this alloy at either annealing temperature. Additions of phosphorus and boron in combination with titanium (MnCrCTiBP) produced a non-uniform distribution of some smaller (80-140 nm) and larger (0.3-1/~m) MC precipi-
tates, both in the matrix and at grain boundaries of the specimens annealed at both temperatures. The specimen annealed at 1050 °C had more dislocations and a higher concentration of annealing twins (Fig. 9) relative to the specimen annealed at 1150 °C. Many of the dislocations were within the planar twin boundaries or extended into the matrix from precipitate particles and appeared to have dissociated into partial dislocations with stacking faults in between. This dislocation configuration suggests that this structure developed during annealing or cooling. None of these features were observed after annealing at 1150°C. "Wavy" large angle grain boundaries were observed after both heat treatments; this morphology appears to be caused by faceting of the grain boundary (Fig. 10(a)). In many instances the facets appeared to be associated with precipitates along the grain boundaries (Fig. 10(b)). Vanadium added in combination with titanium, boron and phosphorus (MnCrCTiVBP) caused a large increase in the amount of fine (15-60 nm) intragranular MC precipitates that formed in the steel annealed at 1050 °C (compared to the steel without vanadium, boron and phosphorus). Many of these MC precipitates appeared in rows or
25
Fig. 11. Microstructure of MnCrCTiVBP annealed 1 h at 1050 °C showing fine MC and coarse M23C6 particles.
Fig. 10. Microstructure of MnCrCTiBP annealed 2 h at 1150 °C showing (a) "wavy" or faceted grain boundary and (b) higher magnification photomicrograph of the precipitates observed in association with wavy boundaries.
bands (Fig. 11 ). There was some grain boundary precipitation--mainly larger (250-800 nm) M23C 6 particles with possibly some smaller (50-100 nm) MC particles in between. After annealing at 1150°C, there was much less precipitation and most of the precipitation was on the grain boundaries; much more grain boundary precipitation was found at 1150 than at 1050 °C. Grain boundary faceting was observed after both annealing treatments. Finally, the addition of titanium, tungsten, vanadium, boron and phosphorus to the base composition (MnCrCTiWVBP) produced a relatively high density of intragranular MC precipitates at 1050 °C, many of them in rows or bands, as observed in the MnCrCTiVBP. As with the MnCrCTiVBP, a much lower density of matrix MC was found after annealing at 1150°C. Many of the particles found in the matrix were surrounded by dislocations. At 1150 °C the few MC precipitates that were observed were on the
grain boundaries, and again grain boundary faceting was evident (Fig. 12). Precipitates noted above were mainly identified in-foil by their characteristic electron diffraction patterns, but some precipitate particles in thin foil specimens annealed at 1150 °C were also analyzed by XEDS. Consistently, most of the coarse particles were chromium-rich M 2 3 C 6 and the finer particles in each sample were MC. Some coarse titanium-rich MC particles were identified in the matrix, but coarse particles at grain boundaries were M23C 6. With the exception of the MnCrC and the MnCrCW, more precipitate formed at 1050 than 1150 °C. This observation is probably due to the difference in carbide solubility at 1050 and 1150 °C. In normal 300 series stainless steels with about 0.05% C, the carbon is fully soluble at 950 and 1050 °C and above. However, in the M n - C r steels with 0.25% C, carbides can form even at 1150°C. The lower carbon solubility at 1050 than at 1150 °C should lead to a larger amount of carbide precipitate forming during the recrystallization and grain growth process at 1050 °C. Since precipitates are observed after the 1050 and 1150 °C anneals, these heat treatments are not true solution anneals.
26
Fig. 12. Grain boundary MC precipitation in the MnCrCTiWVBP annealed 2 h at 1150 °C.
Small TEM disks were used for the microstructural studies. They were cooled rapidly after annealing by pulling them from the furnace into an unheated chamber that was cooled by flowing helium. Such a rapid cooling would be expected to allow them to cool without precipitation occurring. However, the apparent anomalous observation of precipitate--especially grain boundary precipitate--appearing at 1150 and not at 1050 °C for the two alloys without titanium may mean that no precipitate was present when annealed at 1050 and 1150 °C, but precipitation occurred in the steel annealed at 1150 °C because of the slightly longer time that is necessary for a specimen to cool to room temperature from 1150 than from 1050°C. If these precipitates were present at the annealing temperature, an alternative explanation is that the apparent anomaly is due to the dynamics of grain boundary precipitation. That is, after 1 h at 1050 °C, grain boundary motion from recrystallization and/or grain growth processes is continuing, whereas an equilibrium grain size is probably approached in the 2 h at 1150°C. For a system only slightly supersaturated, a precipitate can more easily nucleate on a stationary boundary than a moving one. Matrix precipitates that form in rows or bands after the 1050 °C anneal in several of the other alloys suggest that the precipitation at 1050 °C is intimately connected with the recrystallization process.
3.3. Tensile behavior Figures 13 and 14 show a comparison of the tensile properties of type 316 SS with the Fe-20Mn-12Cr-0.25C base alloy after annealing for 1 h at 1050°C and after being cold worked 20% (common conditions in which such steels are used). The yield stress (YS) of the manganese-stabilized steel in both conditions was equivalent to that of 316 SS (Fig. 13(a)). Since manganese additions cause an increased workhardening rate, the MnCrC steel reached a higher ultimate tensile strength (UTS) when tested in both thermomechanical conditions (Fig. 13(b)). Despite this higher work-hardening capability, the high manganese steel still had equivalent or better ductility (as measured by total elongation) in both the annealed and cold-worked conditions (Fig. 14). These tensile results indicated that an austenitic base alloy with substantial strength and ductility could be obtained by using the information developed from the modified Schaeffier diagram [8]. The next objective was to improve the strength of the new manganese-stabilized alloys. This was attempted by adding titanium, tungsten, vanadium, boron and phosphorus in the quantities given in Table 1 to the F e - 2 0 M n - 1 2 C r 0.25C base composition. Table 4 gives the room temperature tensile properties for the modified alloys, the base alloy and 316 SS as a reference. The manganese-
27 900
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Alloy
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TEMPERATURE (*C) I100
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MnCrC
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4OO I
I
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0
i
i 40o
i
-'~'600
TEMPERATURE (°C)
(b)
Fig. 13. (a) 0.2% yield stress and (b) ultimate tensile strength plotted against test temperature for the F e - M n - C r - C alloy and type 316 stainless steel in the 20% cold-worked condition and annealed 1 h at 1050 °C.
7o
i
I
~a------~. k
~
ANNEALED
0
• MnCrC
-J 50 --
• 316 SS
Ld J
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i
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io o
Elongation (%)
YS
Uniform
Total
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Strength (MPa)
designation -
400 w
TABLE 4 Room temperature tensile properties of manganese-stabilized steels and type 3 1 6 stainless steel
I
4 t
I
20O 400 TEMPERATURE (*C)
I 600
Fig. 14. Total elongation plotted against test temperature for the F e - M n - C r - C (MnCrC) alloy and type 316 stainless steel in the 20% cold-worked condition and annealed 1 h at 1050 °C.
stabilized steels were tested in two annealed conditions (1 h at 1050 °C and 2 h at 1150°C) and in the 20% cold-worked condition. The 316 SS was tested after the 1050 °C solution anneal and after cold working 20%.
Annealed 1 h at 1050 °C MnCrC 220 MnCrCTi 279 MnCrCW 267 MnCrCTiW 302 MnCrCTiBP 288 MnCrCTiVBP 275 MnCrCTiWVBP 304 316 SS 236
798 927 803 918 935 935 915 586
55.4 49.7 57.1 53.8 52.2 51.0 54.9 54.3
56.6 53.0 59.9 56.9 55.6 53.9 57.5 58.2
Annealed 2 h at 1150 °C MnCrC 233 MnCrCTi 258 MnCrCW 247 MnCrCTiW 258 MnCrCTiBP 271 MnCrCTiVBP 221 MnCrCTiWVBP 264
766 891 761 882 891 859 869
53.4 53.5 55.4 54.5 52.8 49.6 59.9
55.1 56.4 57.0 57.2 54.2 50.4 61.7
1086 1160 1057 1168 1158 1126 1114 807
14.1 10.7 17.6 6.6 10.4 11.4 11.3 11.5
16.0 13.0 20.0 9.5 12.1 13.1 13.6 17.4
20% coM worked MnCrC MnCrCTi MnCrCW MnCrCTiW MnCrCTiBP MnCrCTiVBP MnCrCTiWVBP 316 SS
815 954 784 980 946 862 915 739
The room temperature results indicated that the strength of a given high manganese steel annealed for 1 h at 1050 °C generally exceeded that same steel annealed for 2 h at 1150 °C. This behavior probably reflects both more and finer precipitation and a finer grain size for specimens annealed at the lower temperature. With one exception, the YS and UTS of the manganesestabilized stainless steels annealed at 1050°C exceeded those of the 316 SS annealed at this temperature. The exception was that the YS for the MnCrC steel annealed at 1050 °C was slightly lower than the YS of 316 SS. (Note also that only two of the YS values for the manganese-stabilized steels annealed at 1150 °C were less than that of 316 SS annealed at 1050 °C.) The YS and UTS of all of the manganese-stabilized steels in the 20% cold-worked condition exceeded those for 316 SS in that condition. In both annealed conditions the uniform and total elongations of the manganese-stabilized steels at room temperature were similar to those for 316 SS (Table 4). Similar ductility was also observed for most of the alloys in the coldworked condition. The only exception was the MnCrCTiW alloy, which had the lowest uniform
28
and total elongations, although the observed values still indicate adequate ductility. The addition of tungsten (MnCrCW) to the base composition had the least effect on the strength of the annealed steels. An addition of titanium by itself (MnCrCTi) had only a modest strengthening effect. A combination of titanium, boron and phosphorus (MnCrCTiBP) and these three elements plus vanadium (MnCrCTiVBP) gave rise to strengths that were similar to those of MnCrCTi. However, the combination of titanium and tungsten (MnCrCTiW) produced a substantial strength increase over the base' composition. Adding vanadium, boron and phosphorus to the combination of tungsten and titanium did not significantly change the room temperature strength in the annealed condition from that found in MnCrCTiW. In the cold-worked condition tested at room temperature, trends in the relative strength observed for the various combination of elements were similar to those observed in the annealed steels, with the exception that the MnCrCTiW had higher YS and UTS values than the MnCrCTiWVBE However, this greater strength
was also accompanied by lower uniform and total elongation values. Tensile tests were conducted over the range from room temperature to 600 °C. In Figs. 15-20 the tensile properties for the five strongest manganese-stabilized steels (MnCrC and MnCrCW are not shown) are compared with type 316 SS. 60
os4~. ~
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30 0
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~ . ~
\ I 500
600
Fig. 16. Total elongation plotted against test temperature for five heats of manganese-stabilized stainless steels and type 316 stainless steel as a function of test temperature. All steels were annealed 1 h at 1050 °C.
270 320
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~ X
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~
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coo
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S
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450
~ 500
400 400 0 (b)
I 100
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600
Fig. 15. (a) 0.2% yield stress and (b) ultimate tensile strength plotted against test temperature for five heats of manganesestabilized stainless steels and type 316 stainless steel. All steels were annealed 1 h at 1050 °C.
(b)
i
1O0
1
i
i
200 300 400 TEMPERATURE(°C)
i
500
600
Fig. 17. (a) 0.2% yield stress and (b) ultimate tensile strength plotted against test temperature for five heats of manganesestabilized stainless steels and type 316 stainless steel. The manganese steels were annealed 2 h at 1150 °C and the type 316 stainless steel was annealed 1 h at 1050 °C.
29 18
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,
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se
14
i ¸
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i
i
i
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~ ~
52
q
z
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w ~
44
uJ j
8
~
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~
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4
32
0
100
200 300 400 TEMPERATURE (°C)
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600
1000
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~ ~ ~ ~ .-0.,-
MnCrCTi MnCrCTiW MnCrCTiPB MnCrCTiVPB MnCrCTiWVPB
- - C - - 3 1 6 SS
800
00
~0 700 uJ >6O0
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I 100
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1200 11oo
I I I 200 300 400 TEMPERATURE (*C)
~
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~
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~ . , . ~ ~ . ' ~ ~
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I 500
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MnCrCTi MnCrCTiW ~CrCTipB MnCrCTiVPB MnCrCTiWVPB
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i
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I
0
(b)
100
I 200
I I 300 400 TEMPERATUREeC)
0
I 100
I 200
I 300
I 400
I 500
600
TEMPERATURE (°C)
Fig. 18. Total elongation plotted against test temperature for five heats of manganese-stabilized stainless steels and type 316 stainless steel. The manganese steels were annealed 2 h at 1150 °C and the type 316 stainless steel was annealed 1 h at 1050 °C.
.~ " ~ , . ~,,,~'~= . t : ~ ~ ~'~---~ " ",~'111~
2
I 500
600
Fig. 19. (a) 0.2% yield stress and (b) ultimate tensile strength plotted against test temperature for five heats of manganesestabilized stainless steels and type 316 stainless steel. All steels were in the 20% cold-worked condition.
Results clearly indicated a superiority in strength for the manganese-stabilized steels relative to 316 SS after an anneal for 1 h at 1050°C (Figs. 15 and 16), 2 h at 1150 °C (Figs. 17 and 18) and
Fig. 20. Total elongation plotted against test temperature for five heats of manganese-stabilized stainless steels and type 316 stainless steel. All steels were in the 20% cold-worked condition.
after 20% cold work (Figs. 19 and 20). Despite their greater strength, the ductility of the manganese-stabilized steels was generally equivalent to or better than that for 316 SS in the annealed condition (Figs. 16, 18 and 20). These observations on ductility are important because they mean that these new steels represent a significant gain in strength that does not come at the expense of ductility. For the steels annealed for 1 h at 1 0 5 0 ° C the MnCrCTiW and the M n C r C T i W V B P had the highest YS of the modified steels (Fig. 15(a)), with the MnCrCTiWVBP being the strongest at the highest test temperatures. Similar UTS behavior was observed (Fig. 15(b)), with the MnCrCTiWVBP being essentially the strongest alloy over most of the temperature range. Ductilities of all the manganese-stabilized steels annealed at 1050°C were excellent over the entire temperature range and were generally equivalent to or better than those for type 316 SS (Fig. 16). Tensile data for the steels annealed for 2 h at 1150°C (Figs. 17 and 18) displayed the same general behavior as those annealed for 1 h at 1050°C. The MnCrCTiWVBP again had the best strength properties at the highest test temperatures and also had excellent ductility relative to the other steels and type 316 SS. In the cold-worked condition the YS of the MnCrCTiW was highest over most of the temperature range (Fig. 19(a)). The microstructure of this alloy developed stacking faults and dissociated dislocations when aged, suggesting that it has the lowest stacking fault energy of the modified alloys [12], which may contribute to its
30 strength. There was much less difference in the UTS behavior of the five manganese-stabilized steels over this temperature range (Fig. 19(b)). Both YS and UTS of all the modified manganesestabilized steels were substantially greater than found in 20% cold-worked 316 SS. The ductility (Fig. 20) of the MnCrCTiW was lowest and that of the MnCrCTiWVBP was highest for much of the temperature range, although the MnCrCTiWVBP had the lowest ductility at 200 °C. All of the cold-worked F e - M n - C r steels had less ductility than cold-worked 316 SS from room temperature to 200 °C, but then most or all had better ductility from 300 to 600 °C. Fractures for all thermomechanical treatments were ductile and accompanied by necking. The minimum in ductility around 2 0 0 - 4 0 0 °C for the cold-worked, high manganese steels is also observed for cold-worked nickel-containing stainless steels, as seen in Fig. 20 for type 316 stainless steel. Detailed study of this behavior for the fusion prime candidate alloy (PCA), a 16Ni-14Cr-2.5Mo-0.08C austenitic alloy that contains small amounts of titanium, niobium, vanadium, nitrogen and phosphorus, indicated that a cup-cone failure occurred for the low and high temperature tests [13]. A ductile-shear-type fracture with the fracture surface at about 35 ° to the tensile axis occurred at temperatures near the minimum. Scanning electron microscopy studies indicated that this minimum is related to a change from void coalescence at the low and elevated test temperatures to ductile tearing at temperatures of the minimum [13]. These results indicate that it is possible to develop an FIRD steel free of nickel, molybdenum, nitrogen and copper, with tensile properties that are better than those of type 316 stainless steel. In fact, the room temperature YS of the 20% cold-worked MnCrCTiWVBP exceeded that of the PCA [11]. However, the rate at which the YS of the MnCrCTiWVBP decreases with temperature is greater than for the PCA, and above about 450 °C the PCA becomes stronger. The PCA has been under development for several years to optimize its properties and microstructural behavior for strength and irradiation resistance. Further optimization of the properties of the manganese steels should also be possible by adjusting the solute additions and thermomechanlcal treatment to affect the solid solution matrix and the precipitate microstructure.
The tensile results generally indicate the importance of minor alloying additions, e.g. titanium for strength, and synergistic effects make combinations of additives, e.g. titanium and tungsten, more important than single-element additions (Table 4 and Figs. 15-20). Steels containing titanium were generally among the strongest; these were basically the same steels that contained fine MC matrix precipitates observed by T E M in steels annealed at 1050°C in this work or in cold-worked steels aged at 800 °C in previous work [12]. The steel with the combination of vanadium and titanium but without any tungsten (MnCrCTiVBP) may also be of interest. Other work shows the MnCrCTiVBP to be very resistant to recrystallization and the formation of intermetallic phases, e.g. a and Laves, when aged at 800 °C in the cold-worked condition [12]. This steel contained little precipitation after annealing, even though vanadium would be expected to react in conjunction with titanium to form MC. The MnCrCTiVBP was generally the weakest of the steels in the annealed condition, but in the cold-worked condition it was similar to the others at higher temperatures. Tungsten additions to this combination of elements definitely improved the strength in both the annealed and the coldworked conditions. This observation emphasizes the fact that synergistic effects of tungsten and titanium appear to be important. Although the experimental steels discussed in this paper have good strength, they must also be weldable and be compatible with the coolant used in the blanket structure in order to be used for building fusion reactors. Corrosion and weldability studies are in progress. The steels are also being irradiated to determine their resistance to irradiation-induced swelling.
4. Summary and conclusions A series of experimental, high manganese austenitic stainless steels were produced. An F e - 2 0 M n - 1 2 C r - 0 . 2 5 C steel was used as a base alloy composition which was then alloyed for improved strength by adding small amounts of titanium, tungsten, vanadium, boron and phosphorus. Austenite stabilization was achieved with manganese and carbon alone, without the use of high nickel and/or high nitrogen concentrations, as is the case in most high manganese austenitic steels. In tensile tests over the range from room temperature to 600 °C the manganese-stabilized
31
steels had 0.2% yield stress and ultimate tensile strength values that were superior to those of type 316 stainless steel, a commercially available nickel-stabilized austenitic stainless steel. The manganese-stabilized steels were superior in both the solution-annealed and cold-worked conditions. The ductility of the manganese steels was as good as or better than that of the 316 stainless steel for most of the test conditions. Acknowledgments We wish to thank the following people who helped in the completion of this work: N. H. Rouse carried out the tensile tests and prepared the electron microscopy specimens; E. A. Kenik and M. L. Santella reviewed the manuscript; Frances Scarboro prepared the manuscript. This research is sponsored by the Office of Fusion Energy, U.S. Department of Energy, under Contract No. DE-AC05-840R21400 with the Martin Marietta Energy Systems, Inc. References 1 R. L. Klueh and E. E. Bloom, in F. A. Garner, D. S. Gelles and E W. Wiffen (eds.), Optimizing Materials for
2 3 4 5 6 7
8 9 10 11
Nuclear Applications, The Metallurgical Society, Warrendale, PA, 1985, pp. 73-85. J. H. G. Monypenny, Stainless Iron and Steel Vol. 2, Chapman and Hall, London, 1954, p. 261. E. Ruedl, D. Rickerby and T. Sasaki, Fusion Technology, Vol. 2, Pergamon, Oxford, 1984, p. 1029. E. Ruedl and T. Sasaki, J. Nucl. Mater., 122-123 (1984) 794-798. H.R. Brager, F. A. Garner, D. S. Gelles and M. L. Hamilton, J. Nucl. Mater., 133-134 (1985) 907-911. A . H . Bott, F. B. Picketing and G. J. Butterworth, J. NucL Mater., 141-143(1986) 1088-1096. M. Tamura, H. Hayakawa, M. Tanimura, A. Hishinuma and T. Kondo, J. NucL Mater., 141-143 (1986) 1067-1073. R. L. Klueh, P. J. Maziasz and E. H. Lee, Mater. Sci. Eng., A102 (1988) 115-124. P. J. Maziasz and T. K. Roche, J. NucL Mater., 103-104 (1981) 797-802. P. J. Maziasz et aL, J. Nucl. Mater., 108-109 (1982) 296-298. P.J. Maziasz, in B. L. Bramfitt, R. L. Benn, C. R. Brinkman and G. E Vander Voort (eds.), Optimization of Processing, Properties, and Service Performance Through Microstructural Control, American Society for Testing
and Materials, Philadelphia, PA, 1988, pp. 116-161. 12 P.J. Maziasz and R. L. Klueh, in R. L. Klueh, D. S. Gelles, M. Okada and N. H. Packan (eds.), Reduced Activation
Materials for Fusion Reactor Applications, ASTM Spec. Tech. Publ., 1047 (1990) 56-79. 13 D. N. Braski and P. J. Maziasz, J. Nucl. Mater., 122-123 (1984) 338-342.