Effect of heat treatment on microstructure and mechanical properties of ...

Effect of heat treatment on microstructure and mechanical properties of Cr–Ni–Mo–Nb steel M. Q. Wang*, H. Dong, W. J. Hui and J. Shi The microstructure and mechanical properties of a medium carbon Cr–Ni–Mo–Nb steel in quenched and tempered conditions were investigated using transmission electron microscopy (TEM), X-ray analysis, and tensile and impact tests. Results showed that increasing austenitisation temperature gave rise to an increase in the tensile strength due to more complete dissolution of primary carbides during austenitisation at high temperatures. The austenite grains were fine when the austenitisation temperature was ,1373 K owing to the pinning effect of undissolved Nb(C,N) particles. A tensile strength of 1600 MPa was kept at tempering temperatures up to 848 K, while the peak impact toughness was attained at 913 K tempering, as a result of the replacement of coarse Fe rich M3C carbides by fine Mo rich M2C carbides. Austenitisation at 1323 K followed by 913 K tempering could result in a combination of high strength and good toughness for the Cr–Ni–Mo–Nb steel. Keywords: Martensitic steel, Quenching, Tempering, Carbide, Mechanical properties

Introduction Medium carbon low alloy Cr–Ni–Mo–V steels such as G4335V are widely used in quenched and tempered conditions for general engineering and ordnance applications. They also find use in the power generation industry for the production of thick section rotor forgings.1–3 These steels usually contain an amount of Mo element up to 1 mass% so as to avoid temper embrittlement by hindering the segregation of impurity elements such as P to grain boundaries.4 The addition of Mo into the steels can also improve their elevated temperature hardness and strength to some extent due to a secondary hardening effect of fine Mo2C carbides.5 However, a stronger secondary hardening effect is expected for some components subject to severe service environment where the elevated temperature strength is of great importance. In this case, the Mo content should be increased to .1 mass%. It is clear that a very strong secondary hardening effect can be achieved at 3 mass%Mo while a further increase in Mo content cannot evidently increase the hardness.6 Accordingly, a Cr–Ni–Mo–Nb steel containing nominal 3 mass%Mo as well as 3 mass%Cr has been developed recently to have higher elevated temperature strength than the commonly used G4335V steel.7 In the Cr–Ni–Mo–Nb steel, Nb is used instead of V as a microalloying element to prevent abnormal growth of prior austenite grains during austenitisation since Nb is usually found to be more effective in refining austenite National Engineering Research Centre of Advanced Steel Technology, Central Iron and Steel Research Institute, #76 XueYuanNanLu, Beijing 100081, China *Corresponding author,email [email protected]

ß 2007 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 24 February 2006; accepted 11 March 2006 DOI 10.1179/174328407X192796

grains than V during austenitisation at high temperatures.8 Despite these, the microstructure and mechanical properties of the Cr–Ni–Mo–Nb steel have not been fully understood in previous studies. Thus, the present study aims to investigate the effect of austenitisation and tempering on the microstructure and tensile strength and impact toughness of the Cr–Ni–Mo–Nb steel.

Experimental The chemical composition of the steel investigated in the present study is Fe–0?28C–0?13Si–0?20Mn–0?006P– 0?001S–2?99Cr–0?72Ni–2?86Mo–0?11Nb. A high purity melt of the steel was obtained by non-vacuum melting and electroslag remelting (ESR), which was then cast into an ingot of 1000 kg. The ingot was reheated at 1523 K and then forged to round bars of 90 mm in diameter, followed by annealing at 953 K for 24 h to reduce the hardness for easy machining. Round bars of the steel were first austenitised in an air furnace at 1223–1473 K for 1 h, and then quenched into water to room temperature. One piece was kept for each austenitisation temperature and the other pieces were tempered at 473–933 K for 1–10 h. Tensile and Charpy V notch impact specimens were machined from the heat treated bars in the transverse orientation. Finishing machining resulted in tensile specimens 5 mm in diameter and 25 mm in gauge length as well as impact specimens in dimensions of 10610655 mm. Tensile tests were conducted at room temperature at a strain rate of 1022 s21 using an AMSLER-50 material testing machine, and impact tests were performed at room temperature on a standard impact machine. Optical micrographs were obtained by means of the standard metallography using a 3% Nital solution as the etchant. Thin film specimens for transmission electron

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Effect of heat treatment on microstructure and mechanical properties of Cr–Ni–Mo–Nb steel

1 Optical micrographs of specimens austenitised at a 1323 K and b 1473 K for 1 h, water quenched and then tempered at 913 K for 2 h

microscopy (TEM) observations were cut from the heat treated bars, mechanically ground and polished to a thickness of ,50 mm, and then electrochemically thinned with a double jet polishing machine in a 5% perchloric acid (HClO4) and 95% methanol (CH3OH) solution at 233 K. TEM observations were carried out on a JEM-2000FX type microscope, operating at 160 kV using bright field imaging and selected area diffraction (SAD) patterns to identify the carbides. The attached energy dispersive X-ray spectroscopy (EDX) equipment was also used to estimate the chemical composition of the carbides. The fractions of alloy components in carbides were determined through analysing electrolytically extracted carbide residues using X-ray diffraction. Fracture surfaces of impact specimens were observed on the JSM-6400 type scanning electron microscope (SEM), operating at 15 kV.

Results and discussion Effect of austenitisation temperature Figure 1 shows the optical micrographs of the specimens austenitised at 1323 and 1473 K for 1 h, water quenched, and then tempered at 913 K for 2 h. While the 1473 K austenitised specimen had coarse prior austenite grains of .100 mm, the 1323 K austenitised

specimen showed relatively fine austenite grains of ,30 mm, which were similar to those in the 1223 K austenitised specimen. This means that the prior austenite grains did not evidently grow at an austenitisation temperature ,1323 K. In fact, further observations showed that the austenite grains apparently grew only when the austenitisation temperature was .1423 K. Figure 2 shows the TEM images of the as quenched specimen austenitised at 1223 K and then water quenched. Two kinds of particles were observed, one of which was angular with an equivalent diameter of 200–500 nm while the other was spherical with a diameter of ,100 nm. EDX analysis results showed that the large angular particles were rich in Mo and Cr while the small spherical particles were Nb rich. According to the selected area electron diffraction pattern and EDX results, these two kinds of particles were presumably M6C and Nb(C,N) respectively. Figure 3 shows the TEM images of the as quenched specimens austenitised at 1273, 1323 and 1423 K respectively. Compared with Fig. 2, Fig. 3a shows that the M6C particles became spherical and were smaller due to partial dissolution when the austenitisation temperature was 1273 K. When the austenitisation temperature was increased to 1323 K, the M6C particles disappeared and only small Nb(C,N) particles were

2 Bright field images (TEM) of as quenched specimen austenitised at 1223 K and then water quenched, showing a M6C particles and b Nb(C,N) particles respectively

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4 Dependence of mechanical properties on austenitisation temperature for specimens austenitised at various temperatures for 1 h, water quenched and then tempered at 913 K for 2 h

3 Bright field images (TEM) of as quenched specimens austenitised at a 1273 K, b 1323 K and c 1423 K respectively

observed (Fig. 3b). The Nb(C,N) particles did not evidently change when the austenitisation temperature was increased to 1423 K (Fig. 3c). The dependence of mechanical properties of the specimens on the austenitisation temperature is shown in Fig. 4. A tempering temperature of 913 K was used because it would result in high impact toughness, as will be discussed below in the part concerning the effect of tempering temperature. As can be seen from Fig. 4, the tensile strength increased with increasing austenitisation temperature, while the impact toughness exhibited a peak value at the austenitisation temperature of 1373 K.

As shown in Fig. 4, a high austenitisation temperature is expected to improve the tensile strength as a result of the relatively complete dissolution of M6C carbides in the present steel which contained 2?86 mass%Mo as well as 2?99 mass%Cr. However, the austenite grains would grow abnormally when the austenitisation temperature was .1423 K, resulting in a decrease in the impact toughness. Taking into account the nitrogen content of 0?012 mass% in the present Cr–Ni–Mo–Nb steel, the complete dissolution of Nb(C,N) particles could be theoretically calculated to be ,1590 K according to the solubility product proposed by Nordberg and Aronsson.9 It has been suggested that the grain coarsening temperature is ,150 K below the complete dissolution temperature of the Nb(C,N) particles.10 Accordingly, the calculated austenite grain coarsening temperature of the present Cr–Ni–Mo–Nb steel was ,1440 K, which is in good agreement with experimental observations. The complete dissolution of V carbides in two steels containing 0?31C–0?13V and 0?31C–0?34V (mass%) was reported to be 1248 K (Ref. 11). For the two Cr–Ni– Mo–V steels containing 0?35C–0?22V and 0?29C–0?15V (mass%), Ridley and co-workers obtained by theoretical calculations the grain coarsening temperatures of 1025 and 1248 K respectively.2 Nb was used as the microalloying element instead of V in the present Cr–Ni–Mo– Nb steel since the grain coarsening temperatures of the steel were much higher than those of the V microalloyed steels, indicating the effectiveness of Nb as an microalloying element for refining austenite grains during austenitisation at high temperatures up to 1400 K. Thus, an austenitisation temperature of 1323 K was applicable for the Cr–Ni–Mo–Nb steel.

Effect of tempering temperature Observations (TEM) of the specimens tempered at 473 and 573 K showed that tempering at low temperatures gave rise to the formation of e-carbides on the {110} planes of martensitic matrix, and the e-carbides were coarse, which were 100–300 nm in length and 20 nm in thickness when tempered at 573 K. EDX analysis showed that Mo and Cr contents in e-carbides were 2?8 and 3?2 mass% respectively. These results were


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Table 1 Contents of carbides in specimens austenitised at 1323 K for 1 h, water quenched and then tempered at various temperatures for 2 or 10 h, mass% Nb(C,N) M3C M2C M7C3 Total

No. Heat treatment #1 #2 #3 #4 #5 #6 #7

5 Bright field images (TEM) of specimens austenitised at 1323 K for 1 h, water quenched and then tempered for 2 h at various temperatures of a 773 K, b 823 K, c 873 K and d 913 K respectively

similar to the Mo and Cr contents in the steel, indicating that no obvious diffusion of alloy elements occurred during the formation of e-carbides, which was in agreement with the fact that no partitioning of Cr and Mo was observed between carbides and martensite after tempering at 623 K, as reported by Thomson and Miller.12 Figure 5 shows the TEM images for the specimens austenitised at 1323 K for 1 h, water quenched and then tempered at high temperatures for 2 h. In the 773 K tempered specimen, M3C carbides were observed instead of e-carbides, which were shorter but more uniform with the length of ,150 nm, as shown in Fig. 5a. The M3C

As quenched Tempered at 773 Tempered at 823 Tempered at 873 Tempered at 913 Tempered at 913 Tempered at 933

K K K K K K

for for for for for for

0.13 2 h 0.13 2 h 0.13 2 h 0.13 2 h 0.13 10 h 0.13 2 h 0.13

0.18 2.15 2.43 2.45 0.84 1.00 0.94

0.02 0.18 1.03 1.40 2.71 2.89 2.74

0 0 0 0.01 0.10 0.21 0.10

0.33 2.46 3.59 3.99 3.77 4.23 3.91

carbides coarsened up to 600 nm in length when the tempering temperature was increased to 823 K, but the thickness was still ,20 nm, as shown in Fig. 5b. The fraction of M3C carbides also increased when the tempering temperature was increased to 823 K. Some fine carbide less than 50 nm in length could be observed in the 823 K tempered specimens, which were too small to be identified by TEM diffraction. The fine carbides were presumably M2C carbides according to the chemical composition tested by X-ray analysis. The M3C carbides in the 873 K tempered specimen were shorter than those in the 823 K tempered specimen, as shown in Fig. 5c. Tempering at 913 K resulted in a large amount of M2C carbides (Fig. 5d). The M2C carbides could be clearly identified when the tempering time was as long as 10 h or the tempering temperature was increased to 933 K, as shown in Fig. 6. From the SAD patterns, the orientation relationship between M3C carbides and martensite in the present steel was found to obey the Bagaryatski relationship, such that (001)cm//(1¯12)M, [010]cm//[111¯]M and [100]cm// [11¯0]M.13 On the SAD pattern of M2C carbides, marked streaking was usually observed, as shown in Fig. 6. The streaks were reported as a result of the needle morphology of M2C carbides, which leads to discs in reciprocal space and intersects with the Ewald sphere to produce streaks.14 The orientation relationship between M2C carbides and martensite was often assumed to be as follows, (0001)M2C//(110)M and [112¯0]M2C//[100]M. Table 1 lists the mass fractions of various types of carbides in the as-quenched and the quenched and tempered specimens obtained by X-ray analysis on

6 Bright field images (TEM) of specimens austenitised at 1323 K for 1 h, water quenched and then tempered at a 913 K for 10 h and b 933 K for 2 h

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7 Variations of atomic fractions of major alloy elements Fe, Cr and Mo in a M3C carbides and b M2C carbides with tempering temperature for specimens austenitised at 1323 K for 1 h, water quenched and then tempered at various temperatures for 2 h

eletrolytically extracted residues. It can be seen that the content of Nb(C,N) particles was 0?13 mass% in various heat treated conditions and it did not evidently vary with the increase in tempering temperature and tempering time, indicating the high stability of the Nb(C,N) particles during tempering. Actually, the Nb(C,N) particle also included a small amount of Mo element despite the fact that almost all the Nb in the steel was in the form of carbonitride. When the tempering temperature was in the range 773–873 K, the content of M3C carbides was relatively high. Although M2C carbides could not be clearly found by TEM observations in specimens tempered at ,823 K, X-ray analysis proved their existence when the specimen was tempered at .773 K. The content of M2C carbides was relatively high when the tempering temperature was increased to 913 K. An amount of M7C3 carbides was also found in the specimens tempered at high temperatures over 873 K, indicating the tendency of precipitation of equilibrium carbides as were found in Cr–Ni–Mo–V steels.15,16 The total content of carbides increased with increasing tempering temperature except that a slight decrease occurred at ,913 K tempering due to the mass transformation from M3C carbides to M2C carbides. Thus the sequence of carbide precipitation in the present Cr–Ni–Mo–Nb steel during tempering was from ecarbides, to M3C carbides, to M2C carbides and finally to M7C3 carbides. The atomic fractions of major alloying elements in the M3C and M2C carbides are shown in Fig. 7. It can be seen that the M3C carbides contained mainly Fe and a small fraction of Cr and Mo when the tempering temperature was ,873 K. When the tempering temperature reached 913 K, the Cr and Mo fractions in the M3C carbides increased sharply and became comparable with that of Fe. In M2C carbides, the main alloying element was Mo, but the Mo fraction decreased slightly with increasing tempering temperature, indicating that more M2C carbides formed at the expense of Cr from the martensite matrix. It is known that the precipitation of carbides during tempering is strongly related to diffusion of alloying

elements in the martensite matrix. Since the diffusivity of Mo in the martensite matrix is lower than that of Cr, Mo rich M2C carbides likely form at relatively high tempering temperatures, while Fe rich and Cr rich M3C carbides form at relatively low tempering temperatures. Figure 8 shows the distributions of Cr and Mo in the martensite matrix, M3C carbides and M2C carbides in the present Cr–Ni–Mo–Nb steel, which contained totally 2?99 mass%Cr and 2?86 mass%Mo. It can be seen that both Cr and Mo elements were mainly in the martensite matrix at low temperature tempering, while most Mo element was in M2C carbides when the tempering temperature was .873 K, indicating that Mo could only diffuse sufficiently to form M2C carbides at high temperatures. Figure 9 shows the dependence of mechanical properties on tempering temperature for the specimens austenitised at 1323 K for 1 h, water quenched and then tempered at various temperatures for 2 h. When the tempering temperature was ,573 K, increasing tempering temperature resulted in evident decreases in both tensile strength and impact toughness, presumably due to precipitation of M3C carbides and coarsening of e-carbides. When the tempering temperature was in the range 673–848 K, the tensile strength was ,1600 MPa and did not change evidently. Apparently, it was due to secondary hardening effects of carbides that formed during tempering. Owing to the secondary hardening effect of fine M2C carbides, a peak hardness of tempering of Mo containing steels was often observed at ,823 K depending on chemical compositions.17,18 For the present Cr–Ni–Mo– Nb steel, the tensile strength decreased when the tempering temperature reached 873 K. Increasing tempering temperature from 848 to 913 K could result in drastic declines in tensile strength as well as an increase in impact toughness. The impact toughness peaked at a value of 44 J when the tempering temperature was 913 K, followed by a decrease due to coarsening of M2C carbides as the tempering temperature increased to 933 K. Typical fracture surfaces of the Cr–Ni–Mo–Nb steel obtained by SEM observations are shown in Fig. 10.


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8 Variations of mass fractions of major alloy elements a Cr and b Mo in a matrix, M2C and M3C carbides for specimens austenitised at 1323 K for 1 h, water quenched and then tempered at various temperatures for 2 h

9 Dependence of mechanical properties on tempering temperature for specimens austenitised at 1323 K for 1 h, water quenched and then tempered at various temperatures for 2 h

Basically, the fracture at room temperature was ductile in the form of microvoid coalescence, despite the fact that the tensile strength was different and the impact toughness varied with tempering temperature. The size of dimples in the 913 K tempered specimen, which had the peak impact toughness, was much larger than that in the 823 K tempered specimen, as can be seen from Fig. 10. Accordingly, the heat treatment rout, in which the steel was austenitised at 1323 K for 1 h, water quenched and then tempered at 913 K for 2 h, should result in a good combination of high strength and good toughness for the Cr–Ni–Mo–Nb steel. After heat treated in such a route, the present Cr–Ni– Mo–Nb steel also exhibited higher elevated temperature strengths than the G4335V steel. It showed a 12% loss of its room temperature strength when tested at 623 K while that of the G4335V was .20% (Ref. 7). In comparison with the G4335V steel which should have a small amount of M2C carbides when tempered at high temperatures, the present Cr–Ni–Mo–Nb steel showed relatively high amount of M2C carbides, which should contribute to its high elevated temperature strength due

10 Typical fracture surfaces of impact specimens austenitised at 1323 K for 1 h, water quenched and then tempered at a 823 K and b 913 K for 2 h

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to the strong secondary hardening effect of M2C carbides.

Conclusions The dissolution of primary carbides during austenitisation and the precipitation of carbides during tempering for the Cr–Ni–Mo–Nb steel were investigated using TEM and X-ray analysis. The mechanical properties in various heat treatment conditions were also obtained by tensile and Charpy impact tests. The following conclusions can be drawn. 1. The primary M6C carbides could dissolve completely only when the austenitisation temperature was as high as 1323 K due to the high Mo content in the Cr– Ni–Mo–Nb steel. The prior austenite grains did not evidently grow when the austenitisation temperature was ,1423 K owing to the pinning effect of Nb(C,N) particles. 2. With increasing austenitisation temperature from 1223 to 1473 K, the tensile strength increased and the impact toughness peaked at 1373 K austenitisation due to the more complete dissolution of primary carbides and the evident growth of prior austenite grains at 1423 K. 3. With the increase in tempering temperature, e-, M3C, M2C and M7C3 carbides successively precipitated from the martensitic matrix. X-ray analysis results showed that the M3C and M2C carbides were Fe rich and Mo rich respectively. At high temperature tempering, the Mo element was mainly in the M2C carbides while the Cr element was mainly in the martensite matrix.

4. The Cr–Ni–Mo–Nb steel had a strong secondary hardening effect during tempering at ,828 K as a result of the mass precipitation of fine M2C carbides. The 913 K tempered specimen showed the highest impact toughness.

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