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Hot workability of spray-formed AISI M3:2 high-speed steel
Materials Science and Engineering A 386 (2004) 420–427
Hot workability of spray-formed AISI M3:2 high-speed steel C. Rodenburg, M. Krzyzanowski, J.H. Beynon, W.M. Rainforth∗ IMMPETUS, Department of Engineering Materials, University of Sheffield, Sir Robert Hadfield Building, Mappin Street, Sheffield S1 3JD, UK Received 19 December 2003; received in revised form 16 July 2004
Abstract Axisymmetric hot compression tests (900–1100 ◦ C) on spray-formed AISI M3:2 high-speed steel were performed in order to establish suitable parameters for hot forging of this material. Special attention was paid to establish the deformation conditions that lead to the breakdown of the carbide network, present after spray forming, and to avoid fracture of the material as a result of deformation. By a combination of microstructural analysis and finite element modelling, values for fracture stresses in this temperature range and critical strains for the breakdown of the carbide network are given. The activation energy for hot deformation was also determined. © 2004 Elsevier B.V. All rights reserved. Keywords: High-speed steel; Hot compression; Activation energy; Carbide network; Fracture stress
1. Introduction In the spray-forming process the metal melt is ‘atomised’ in a spray chamber in a protective gas atmosphere and the spray is collected on a rotating former. By simultaneous rotation and translation of the former parallel to the spray axis, the semi-liquid deposit builds up to form a solid of the required shape [1]. Consequently, cooling rates involved in spray forming differ from those found in casting or powder metallurgical methods. Therefore, spray-formed materials differ significantly in their microstructure from cast materials or even powder metallurgical manufactured materials [2]. For example, spray-forming offers a reduction in microstructural scale and a reduction in the size and interconnectivity of carbides in hot work tool steels compared with the conventionally cast equivalent [3]. Although it is now possible to produce high-density spray-formed steels in large sizes, some residual porosity is still found. Therefore, high-stress applications such as work rolls for hot rolling mills require hot working of the spray-formed steel. There is little work published on the hot workability of tools steels in general [4–8] but in particular the hot work∗
Corresponding author. Tel.: +44 114 2225469; fax: +44 114 2225943. E-mail address: [email protected] (W.M. Rainforth).
0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.07.056
ability of spray-formed tools steels has not been addressed, though there is limited work on the effect of hot working conditions on carbide size and distribution [2]. The different behaviour of spray-formed material and cast material was shown for 17% Cr white cast iron, which was hot forged in the spray-formed condition, hitherto impossible for such a material [9]. In this paper, the hot workability of spray-formed AISI M3:2 high-speed steel is investigated by means of hot compression testing in an axisymmetric geometry. The experimental results are interpreted with the help of finite element (FE) modelling in order to account for strain variations within the specimen. 2. Experimental techniques A 3-ton spray-formed high-speed steel billet was supplied by Danspray (Gregersenvej 8, 2630 Taastrup, Denmark) in the as-sprayed condition with approximate dimensions 0.5 wide by 2 m long. The composition, which conforms to the ASP2023 grade (usually used for powder processing) of the spray-formed high-speed steel, was measured as (in weight %): 1.31 C, 6.24 W, 5.34 Mo, 4.12 Cr, 3.05 V, 0.56 Si, 0.24 Mn, 0.11Ni, 0.22 Cu, and 749 ppm N, remainder Fe.
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Table 1 Dimensions of specimens before and after deformation Nr.
Tnominal (◦ C)
Texperiment (◦ C)
Strain rate (s−1 )
d0 (mm)
h0 (mm)
df (mm)
hf (mm)
B
A B C D E F G H I
900 900 900 1000 1000 1000 1100 1100 1100
885 886 889 973 981 987 1075 1063 1080
0.1 1 10 0.1 1 10 0.1 1 10
11.99 11.97 12.02 11.97 12.03 11.99 12.02 11.97 12.02
14.97 14.93 15.00 14.94 14.92 15.01 14.91 14.98 14.95
17.77 17.50 17.73 17.71 17.71 17.76 17.58 17.65 17.72
7.73 7.80 7.76 7.69 7.78 7.85 7.77 7.70 7.74
1.13 1.12 1.13 1.13 1.13 1.15 1.11 1.12 1.13
d0 , h0 , df , hf and B are defined in the text.
Because it is well known that despite having better homogeneity than a conventionally cast material, there is nevertheless some microstructural inhomogeneity across the diameter of a spray-formed billet [10]; a microstructural inspection across the billet diameter was performed first to ensure that the initial microstructure for all compression experiments would be identical. A combination of hardness testing and light microscopy was used to identify a reasonably homogeneous region. On the basis of the microstructural study, cylindrical test specimens with an initial diameter d0 ≈ 12 mm and of initial height h0 ≈ 15 mm were removed from the spray-formed billet. The orientation of the samples with respect to the original billet was the same throughout. Axisymmetric compression tests [11] at temperatures in the range of 900–1100 ◦ C and strain rates between 0.1 and 10 s−1 were carried out in order to simulate the microstructure evolution during forging. The nominal maximum strain for all experiments was 0.69, which corresponds to a reduction of 50%. This is a typical height reduction for billets to be forged into work rolls. The load was applied by means of ceramic platens with an average temperature, as given in Table 1. There was a temperature gradient between upper and lower platen e.g. for a nominal test temperature of 1000 ◦ C, the temperatures of the platens were 990 and 1010 ◦ C, respectively. The temperature of the test specimen was measured by means of a thermocouple in the centre of the specimen. The specimen was heated by induction for 2 min to the required test temperature and then held at that temperature for 1 min before the compression test was started. After reduction of the specimen to 50% of the original height, the test specimen was removed from the furnace and water quenched. The time between the end of deformation and onset of water quench was about 1 s.
3. Finite element modelling There are significant differences in deformation developed in different regions within the sample during uni-axial compression testing. A non-steady-state mathematical model
based on application of the FE method was used to calculate the strain and the strain rates in the local areas for analysis of observed microstructure variations across the specimen during the testing. For the purpose of the analysis, the calculations were made only within the quarter of the specimen cross-section by assuming symmetry along the main axis of the specimen. The FE model consisted of 1016 four-node, isoparametric, arbitrary quadrilateral axisymmetric elements forming the tool and the specimen. Normally, the compression tests were carried out using position control. Thus, an application of the corresponding displacements at the tool edge simulated loading in the model. The specimen was assumed to be elastic–plastic while the tool was simulated as an elastic body. The mechanical properties of the tool and the specimen were introduced on the basis of available experimental data [12]. The temperature-corrected flow curves obtained during the compression testing of the material were used for prediction of the specimen’s plastic deformation. The sliding between the tool and the specimen surfaces during the compression testing was modelled using a shearbased model of friction such that v 2 rel ¯t τ = −m kY arctan (1) π c where τ is the shear stress; m is the friction factor, here assumed to be 0.35; kY is the shear yield stress; c is a constant taken to be about 1% of a typical relative velocity vrel , which smoothes the discontinuity in the value of the shear stress when stick/slip transfer occurs between the surfaces in a contact; and ¯t is the tangent unit vector in the direction of the relative sliding velocity. The commercial MARC K7.2 finite element code was used to calculate the metal flow and the frictional sliding during the compression of the axisymmetrical specimens. The tangential separation stresses were calculated in the model using the deformable–deformable contact procedure implemented in the MARC code. Once contact between a node and a deformable surface is detected, a tie is activated. The tying matrix was such that the contacting node could slide along
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the surface, be separated or be stuck, according to the general contact conditions.
4. Results 4.1. Dimensions and appearance of compressed specimens The dimension of the samples before and after compression are summarised in Table 1. Here d0 and h0 are the initial diameter and the initial height and df and hf are the deformed diameter and height. These dimensions are used to judge the validity of the test data and also serve as parameters to adjust the FE model. A validity criterion for data obtained by the axisymmetric compression test is the barrelling coefficient B [11] B=
hf df2 . h0 d02
(2)
Tests for which B ≥ 1.1 are invalid for the purpose of flow stress determination [11]. This criterion, applied to the tests summarised in Table 1, renders most tests invalid as far as the exact determination of the flow stress is concerned. The large values of B indicate that the data should be friction corrected. However, the small number of tests does not provide sufficient information to extract friction data to perform an appropriate friction correction. Therefore, instead of stress–strain curves, mean pressure strain curves are given. Specimens deformed at 900 ◦ C at all the strain rates exhibited cracks along the circumference with the crack running in the direction of the compression axis. 4.2. Flow curves The recorded raw data were converted into mean pressure strain curves using the following relations. The strain ε is given by ε = ln
h , h0
Fig. 1. Flow curves for all experiments. The dotted line indicates the positions of peak mean pressure values for the determination of the activation energy.
presented in Fig. 2. As expected, the largest effect is observed for the highest strain rate. The flow curves can be temperature corrected if the activation energy for hot working QHW is known. 4.3. Determination of activation energy The activation energy for hot working was calculated using the method described in [13], which assumes that the influence of temperature T and strain rate ε˙ on flow stress σ is given by QHW A exp(βσ) = ε˙ exp , (5) RT where A and β are constants and R is the universal gas constant. Mean pressure values were determined by taking the intersection of the flow curve and the dotted line in Fig. 1. With these values an initial activation energy of 505 kJ mol−1 was estimated. However, as can be seen by comparing Figs. 1 and 2, these peak values are already affected by significant deformation heating. Therefore, the initial activation energy of 505 kJ mol−1 was used to correct the ex-
(3)
where h is the height during each stage of the test and h0 is the initial height. The mean pressure P is given by 4F , (4) πd 2 where F is the force at any given time and d is the corresponding specimen diameter at that time. The mean pressure strain curves obtained in this way are shown in Fig. 1. Although all curves exhibit a clear maximum mean pressure, it is debatable if this is a result of dynamic recrystallisation or deformation heating. From the recorded temperature data it is evident that deformation heating takes place, as can be seen from typical strain temperature curves for a nominal deformation temperature of 900 ◦ C, which are P=
Fig. 2. Strain temperature curves for a nominal deformation temperature 900 ◦ C and three strain rates.
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elling was also used to determine the hoop stress experienced for different test conditions, since it was found that the hoop stress for deformation at 900 ◦ C exceeded the fracture stress of the material, as evident by the occurrence of cracks. The
Fig. 3. Temperature-corrected flow curves. The dotted line indicates the positions of mean pressure values for the determination of the activation energy in last iteration. Experiment D is not shown due to the loss of thermocouple in the early stages of experiment. Without a temperature recording during the test, a temperature correction could not be performed.
perimental flow curves for deformation heating. From these temperature-corrected flow curves, a new activation energy was derived. The iteration process was repeated until no change in activation energy within the error boundaries was observed. The error was defined as the variation of activation energy in the flow stress region investigated here, which covers the range from 150 to 450 MPa. The final value for the activation energy by this method was found to be 487 ± 31 MPa. The temperature-corrected flow curves using this final activation energy are presented in Fig. 3. 4.4. Strain distribution obtained from FE model An example for the strain distribution obtained from the FE model is shown in Fig. 4, which represents deformation at a nominal temperature of 900 ◦ C and a strain rate of 10 s−1 . The positions labelled in this figure will be explained below in conjunction with microstructural investigations. FE mod-
Fig. 4. Strain distribution obtained from FE model for experiment C. Labels (a)–(d) mark positions, from which optical micrographs in Fig. 6(a)–(d) were taken.
Fig. 5. Microstructure of as-sprayed material (a) optical micrograph of polished sample demonstrating the presence of a carbide network, (b) optical micrograph of etched sample, (c) SEM image of etched sample.
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C. Rodenburg et al. / Materials Science and Engineering A 386 (2004) 420–427 Table 2 Summary of critical strain and strain rate at position where carbide network was broken down completely, obtained from FE modelling Experiment
Tnominal (◦ C)
εcrit
ε˙ (s−1 )
A C G
900 900 1100
0.88 0.89 0.82
0.18 16.8 0.17
minimum hoop stress for deformation at 900 ◦ C was obtained for a strain rate of 0.1 s−1 and was estimated to be 1.8 GPa. For deformation conditions resulting in smaller hoop stresses (i.e. at higher deformation temperatures), no cracks were observed. However, for deformation at 1000 ◦ C and strain rate of 10 s−1 with an estimated maximum hoop stress of 1.9 GPa, no cracks were observed. Therefore, the fracture stress of the investigated high-speed steel is assumed to be in the range 1.8–1.9 GPa over the given temperature range.
Fig. 6. Series of optical micrographs along compression axis of specimen deformed in experiment C. Positions corresponding to (a)–(d) are indicated in Fig. 4.
Fig. 7. SEM images of cross-sections perpendicular to compression axis of compression test samples (a) experiment B: nominal deformation temperature 900 ◦ C, (b) experiment H: nominal deformation temperature 1100 ◦ C. Note the difference in density of submicron-sized carbides and the presence of cracked carbides in (a).
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4.5. Influence of compression test parameters on microstructure The initial microstructure of the as-sprayed material is shown in Fig. 5. As far as the carbide distribution is concerned three different length scales were present: I. carbide network around 70 m consisting of irregular shaped carbides (highlighted by a black line in Fig. 5a), and close to elliptical or round carbides with average size of about 10 microns (marked by white lines in Fig. 5a); II. close to elliptical or round carbides of micron size (examples marked by white circles in Fig. 5c); III. rod shaped and rectangular carbides of sub micrometer size (Fig. 5c). Carbides of category I and II were visible by optical microscopy of polished samples (Fig. 5a). Carbides of group III were only visible in SEM images (Fig. 5c). However, a darkening in optical micrographs of the nital etched material (Fig. 5b) is indicative of their presence. This can be seen by comparing Fig. 5b and c. From these figures it is also evident that the carbide density varies from cell to cell. In order to investigate the influence of compression test parameters on microstructure, these different length scales have to be considered separately. The influence of strain on the carbide network structure is shown in Fig. 6, which shows a series of optical micrographs taken along the compression axis of the test specimen, as indicated in Fig. 4. No deformation of the cells is observed in Fig. 6a, taken from the area in contact with the tool during compression. Moving towards the centre of the specimen the carbide cell is first seen to be deformed in Fig. 6b, and then only local remnants are observed in Fig. 6c. There were no indications for the presence of a carbide network in the centre of the specimen (Fig. 6d). This behaviour was observed for all tests and indicates that the carbide network is broken down by deformation amounting to an equivalent plastic strain between 0.82 and 0.89, depending on deformation temperature. Table 2 summarises values obtained from FE modelling for the critical strain εcrit and strain rates ε˙ at which no carbide network remains. A strong influence of strain rate was not observed. A higher deformation temperature resulted in a slightly lower critical strain. The deformation temperature influences the carbide network in the direction perpendicular to the compression axis. The cracking of large irregular-shaped carbides was observed for deformation at 900 ◦ C (Fig. 7a) but not at 1100 ◦ C (Fig. 7b). Carbides falling into category II are not influenced by high stresses occurring at low deformation temperature, since their round shape and smaller size make them less prone to cracking [14]. Furthermore, these particles were identified by electron beam microanalysis as VC-type carbo-nitrides, which do not dissolve at the temperatures involved. The number of carbides of category III strongly depends on the deformation temperature, as is evident by comparing Fig. 7a and b. In these figures, the small carbides are rep-
Fig. 8. Optical micrographs for specimens deformed at strain rates given and nominal temperatures T = 900 ◦ C (a), T = 1000 ◦ C (b) and T = 1100 ◦ C (c).
resented by the small white dots in the low-magnification SEM images. The low magnification was chosen to provide a better impression of the density distribution over a large area, whereas the insets demonstrate the actual shape and size
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Table 3 Summary of activation energies for hot working a number of tools steels Steel
Q in kJ mol−1
Composition (wt.%) C
Cr
Mo
W
V
T1
0.95
4.35
0.11
18.41
0.99
AISI M3:2 M2 D2 H13 A2 W1
1.31 0.84 1.45 0.38 1.00 1.03
4.12 4.00 11.09 5.25 5.00 0.07
5.34 5.00 0.72 1.30 1.15 0.03
6.24 6.50 – – – –
3.05 1.90 0.71 1.05 0.3 –
of the type III carbides. The high density of carbides at the lower deformation temperature of 900 ◦ C in Fig. 7a is apparent. A low carbide density in conjunction with a strong local density variation was observed for deformation at 1100 ◦ C (Fig. 7b). The deformation temperature also influences the structure of the matrix, as can be seen from Fig. 8. For deformation below 1100 ◦ C (Fig. 8a and b), the matrix grain boundaries were not revealed by nital etching. However, following deformation at 1100 ◦ C, a well-developed matrix grain structure, resembling a recrystallised structure, was observed (Fig. 8c). However, recrystallisation did not seem to have taken place homogenously, and was only observed in the bands of lower carbide density (represented by absence of small black dots) in Fig. 8c.
5. Discussion The activation energy for hot working of 487± 31 kJ mol−1 is in the upper range of activation energies found for other tool steels, as can be seen from Table 3. However, it is conceivable that this is due to the high carbon content in combination with a very high vanadium content in the form of vanadium carbo-nitrides. A change in activation energy due to precipitation, as described in [5], was not observed in the temperature range tested here. The temperature-corrected flow curves for T < 1100 ◦ C do not exhibit maxima that would indicate the presence of dynamic recrystallisation, although dynamic recrystallisation is found to take place in other tool steels at this temperature and strain range for comparable strain rates [5,6]. The corresponding microstructures for T < 1100 ◦ C do not show signs of recrystallisation. In contrast, deformation at 1100 ◦ C does result in a recrystallised microstructure even though there is no pronounced peak in the flow curve. Therefore, static recrystallisation may have taken place during the time (about 1 s) between the end of deformation and the onset of quenching. Furthermore, within the high strain region recrystallisation seems to take place inhomogenously, confined to areas with a low density of submicron carbides. This is consis-
654 467 487 455 428 401 399 286
Ref. <1000 ◦ C >1000 ◦ C
[5] Current work [15] [16] [4] [15] [16]
tent with the presence of closely spaced carbides pinning the grain boundaries (Zener drag mechanism) and hindering recrystallisation [17]. The same mechanism is thought to be responsible for the absence of recrystallisation at deformation at 900 and 1000 ◦ C. As shown in Fig. 7, the carbide density of submicron-sized carbides is much higher at the lower deformation temperature. This higher volume fraction seems to be sufficient to prevent recrystallisation. One of the main aims of hot working of the investigated high-speed steel is the elimination of the carbide network. It was shown with the example of M2 that the dominant mechanism for breaking up the carbide network is mechanical fragmentation as opposed to a diffusion-controlled mechanism [7]. This is supported by this work, where Fig. 6 demonstrates progressive fragmentation with increasing strain. However, a small influence of deformation temperature on the critical strain (0.89 ≥ εcrit ≥ 0.82) was found (see Table 2), indicating that diffusion is involved to a small extent. In the current work, the critical strain for fragmentation of the carbide network was determined for a single deformation. However, in the normal forging of billets, the forging process would be undertaken in progressive smaller strain steps. Nevertheless, the total strain would allow not only break-up of the carbide structure but also static recrystallisation to refine the matrix structure and relax the residual strain in between deformations. The formation of cracks during hot working is undesirable. This work demonstrates that the hoop stress during deformation at 900 ◦ C exceeded the fracture stress and, therefore, should be avoided. Applying the values determined in these experiments to industrial practice would require a FE model of the particular industrial forming operation, highlighting the surface stress maxima.
6. Conclusions The hot workability of spray formed AISI M3:2 type high-speed steel was investigated in order to specify suitable forging conditions for this steel. From Arrhenius plots an activation energy for hot working of 487 ± 31 kJ mol−1 was determined. The combination of FE modelling and the
C. Rodenburg et al. / Materials Science and Engineering A 386 (2004) 420–427
observation of cracks for deformation at 900 ◦ C, but not at 1000 ◦ C, led to an estimated fracture stress in the range of 1.8–1.9 GPa. By observing microstructural changes as a function of specimen position and linking this to the strain distribution obtained by FE modelling, a critical strain, which is required to break up the carbide network, typically present in tool steels, was estimated to be 0.89 ≥ εcrit ≥ 0.82 covering the temperature range of 900–1100 ◦ C. This critical strain value has practical implications, for example the forging of spray formed high-speed steel billets into work rolls for hot rolling mills. It was found previously that the presence of the carbide network is responsible for increased oxidative wear at elevated temperatures [18]. Hence, the wear should be reduced if the forging conditions are tailored to exceed the critical strain obtained in this work.
Acknowledgements The authors would like to thank A. Lacey for performing the compression tests and E. Condliffe, School of Earth Sciences, University of Leeds for the electron microprobe analysis. This work was funded by EPSRC under grant number GR/R80209/01. The assistance of Danspray and Sheffield Forgemasters Rolls is also gratefully acknowledged.
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References [1] C. Spiegelhauer, Steel World 5 (2000) 20. [2] R.A. Mesquita, C.A. Barbosa, Proceedings SDMA, vol. 1, Bremen, June 22–25, 2003, ISBN 3-8330-0571-8. [3] P.S. Grant, Prog. Mater. Sci. 39 (1995) 497. [4] C. Imbert, N.D. Ryan, H.J. Mc Queen, Metall. Trans. A 15 (1984) 1855. [5] J. Liu, H. Chang, R. Wu, T.Y. Hsu, X. Ruan, Mater. Charact. 45 (2000) 175. [6] R. Milovi´c, D. Manojloi´c, M. Andjeli´c, D. Drobnjak, Steel Res. 63 (1992) 78. [7] M.R. Ghomashchi, C.M. Sellars, Metall. Trans. A 24 (1993) 2171. [8] J.W. Wang, Z.K. Yao, Iron Steel (China) 30 (1995) 52. [9] D.N. Hanlon, Y.H. Li, W.M. Rainforth, C.M. Sellars, J. Mater. Sci. Lett. 17 (1998) 1637. [10] P. Kjeldsteen, C. Spiegelhauer, O. Sandberg, Proceedings SDMA, vol. 1, Bremen, June 22–25, 2003, ISBN 3-8330-0571-8. [11] B. Roebuck, J.D. Lord, M. Brooks, M.S. Loveday, C.M. Sellars, R.W. Evans, Measuring, Flow Stress in Hot Axisymmetric Compression Tests, National Physics Laboratory, 2002. [12] M. Pietrzyk, J.G. Lenard, Thermal–Mechanical Modelling of the Flat Rolling Process, first ed., Springer, Heidelberg, Berlin, 1991. [13] G.J. Richardson, D.N. Hawkins, C.M. Sellars, Worked Examples in Metalworking, The Institute of Metals, London, 1985. [14] T. Antretter, F.D. Fischer, Mater. Sci. Eng. A 237 (1997) 6. [15] C.A.C. Imbert, H.J. McQueen, Mater. Sci. Eng. A 313 (2001) 88. [16] C.A. Imbert, H.J. McQueen, Mater. Sci. Technol. 16 (2000) 514. [17] F.J. Humphreys, M. Hatherley, Recrystallization and Related Annealing Phenomena, Pergamon, 2002, p. 73. [18] D.N. Hanlon, W.M. Rainforth, Wear 255 (2003) 956.
Hot workability of spray-formed AISI M3:2 high-speed steel C. Rodenburg, M. Krzyzanowski, J.H. Beynon, W.M. Rainforth∗ IMMPETUS, Department of Engineering Materials, University of Sheffield, Sir Robert Hadfield Building, Mappin Street, Sheffield S1 3JD, UK Received 19 December 2003; received in revised form 16 July 2004
Abstract Axisymmetric hot compression tests (900–1100 ◦ C) on spray-formed AISI M3:2 high-speed steel were performed in order to establish suitable parameters for hot forging of this material. Special attention was paid to establish the deformation conditions that lead to the breakdown of the carbide network, present after spray forming, and to avoid fracture of the material as a result of deformation. By a combination of microstructural analysis and finite element modelling, values for fracture stresses in this temperature range and critical strains for the breakdown of the carbide network are given. The activation energy for hot deformation was also determined. © 2004 Elsevier B.V. All rights reserved. Keywords: High-speed steel; Hot compression; Activation energy; Carbide network; Fracture stress
1. Introduction In the spray-forming process the metal melt is ‘atomised’ in a spray chamber in a protective gas atmosphere and the spray is collected on a rotating former. By simultaneous rotation and translation of the former parallel to the spray axis, the semi-liquid deposit builds up to form a solid of the required shape [1]. Consequently, cooling rates involved in spray forming differ from those found in casting or powder metallurgical methods. Therefore, spray-formed materials differ significantly in their microstructure from cast materials or even powder metallurgical manufactured materials [2]. For example, spray-forming offers a reduction in microstructural scale and a reduction in the size and interconnectivity of carbides in hot work tool steels compared with the conventionally cast equivalent [3]. Although it is now possible to produce high-density spray-formed steels in large sizes, some residual porosity is still found. Therefore, high-stress applications such as work rolls for hot rolling mills require hot working of the spray-formed steel. There is little work published on the hot workability of tools steels in general [4–8] but in particular the hot work∗
Corresponding author. Tel.: +44 114 2225469; fax: +44 114 2225943. E-mail address: [email protected] (W.M. Rainforth).
0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.07.056
ability of spray-formed tools steels has not been addressed, though there is limited work on the effect of hot working conditions on carbide size and distribution [2]. The different behaviour of spray-formed material and cast material was shown for 17% Cr white cast iron, which was hot forged in the spray-formed condition, hitherto impossible for such a material [9]. In this paper, the hot workability of spray-formed AISI M3:2 high-speed steel is investigated by means of hot compression testing in an axisymmetric geometry. The experimental results are interpreted with the help of finite element (FE) modelling in order to account for strain variations within the specimen. 2. Experimental techniques A 3-ton spray-formed high-speed steel billet was supplied by Danspray (Gregersenvej 8, 2630 Taastrup, Denmark) in the as-sprayed condition with approximate dimensions 0.5 wide by 2 m long. The composition, which conforms to the ASP2023 grade (usually used for powder processing) of the spray-formed high-speed steel, was measured as (in weight %): 1.31 C, 6.24 W, 5.34 Mo, 4.12 Cr, 3.05 V, 0.56 Si, 0.24 Mn, 0.11Ni, 0.22 Cu, and 749 ppm N, remainder Fe.
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Table 1 Dimensions of specimens before and after deformation Nr.
Tnominal (◦ C)
Texperiment (◦ C)
Strain rate (s−1 )
d0 (mm)
h0 (mm)
df (mm)
hf (mm)
B
A B C D E F G H I
900 900 900 1000 1000 1000 1100 1100 1100
885 886 889 973 981 987 1075 1063 1080
0.1 1 10 0.1 1 10 0.1 1 10
11.99 11.97 12.02 11.97 12.03 11.99 12.02 11.97 12.02
14.97 14.93 15.00 14.94 14.92 15.01 14.91 14.98 14.95
17.77 17.50 17.73 17.71 17.71 17.76 17.58 17.65 17.72
7.73 7.80 7.76 7.69 7.78 7.85 7.77 7.70 7.74
1.13 1.12 1.13 1.13 1.13 1.15 1.11 1.12 1.13
d0 , h0 , df , hf and B are defined in the text.
Because it is well known that despite having better homogeneity than a conventionally cast material, there is nevertheless some microstructural inhomogeneity across the diameter of a spray-formed billet [10]; a microstructural inspection across the billet diameter was performed first to ensure that the initial microstructure for all compression experiments would be identical. A combination of hardness testing and light microscopy was used to identify a reasonably homogeneous region. On the basis of the microstructural study, cylindrical test specimens with an initial diameter d0 ≈ 12 mm and of initial height h0 ≈ 15 mm were removed from the spray-formed billet. The orientation of the samples with respect to the original billet was the same throughout. Axisymmetric compression tests [11] at temperatures in the range of 900–1100 ◦ C and strain rates between 0.1 and 10 s−1 were carried out in order to simulate the microstructure evolution during forging. The nominal maximum strain for all experiments was 0.69, which corresponds to a reduction of 50%. This is a typical height reduction for billets to be forged into work rolls. The load was applied by means of ceramic platens with an average temperature, as given in Table 1. There was a temperature gradient between upper and lower platen e.g. for a nominal test temperature of 1000 ◦ C, the temperatures of the platens were 990 and 1010 ◦ C, respectively. The temperature of the test specimen was measured by means of a thermocouple in the centre of the specimen. The specimen was heated by induction for 2 min to the required test temperature and then held at that temperature for 1 min before the compression test was started. After reduction of the specimen to 50% of the original height, the test specimen was removed from the furnace and water quenched. The time between the end of deformation and onset of water quench was about 1 s.
3. Finite element modelling There are significant differences in deformation developed in different regions within the sample during uni-axial compression testing. A non-steady-state mathematical model
based on application of the FE method was used to calculate the strain and the strain rates in the local areas for analysis of observed microstructure variations across the specimen during the testing. For the purpose of the analysis, the calculations were made only within the quarter of the specimen cross-section by assuming symmetry along the main axis of the specimen. The FE model consisted of 1016 four-node, isoparametric, arbitrary quadrilateral axisymmetric elements forming the tool and the specimen. Normally, the compression tests were carried out using position control. Thus, an application of the corresponding displacements at the tool edge simulated loading in the model. The specimen was assumed to be elastic–plastic while the tool was simulated as an elastic body. The mechanical properties of the tool and the specimen were introduced on the basis of available experimental data [12]. The temperature-corrected flow curves obtained during the compression testing of the material were used for prediction of the specimen’s plastic deformation. The sliding between the tool and the specimen surfaces during the compression testing was modelled using a shearbased model of friction such that v 2 rel ¯t τ = −m kY arctan (1) π c where τ is the shear stress; m is the friction factor, here assumed to be 0.35; kY is the shear yield stress; c is a constant taken to be about 1% of a typical relative velocity vrel , which smoothes the discontinuity in the value of the shear stress when stick/slip transfer occurs between the surfaces in a contact; and ¯t is the tangent unit vector in the direction of the relative sliding velocity. The commercial MARC K7.2 finite element code was used to calculate the metal flow and the frictional sliding during the compression of the axisymmetrical specimens. The tangential separation stresses were calculated in the model using the deformable–deformable contact procedure implemented in the MARC code. Once contact between a node and a deformable surface is detected, a tie is activated. The tying matrix was such that the contacting node could slide along
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the surface, be separated or be stuck, according to the general contact conditions.
4. Results 4.1. Dimensions and appearance of compressed specimens The dimension of the samples before and after compression are summarised in Table 1. Here d0 and h0 are the initial diameter and the initial height and df and hf are the deformed diameter and height. These dimensions are used to judge the validity of the test data and also serve as parameters to adjust the FE model. A validity criterion for data obtained by the axisymmetric compression test is the barrelling coefficient B [11] B=
hf df2 . h0 d02
(2)
Tests for which B ≥ 1.1 are invalid for the purpose of flow stress determination [11]. This criterion, applied to the tests summarised in Table 1, renders most tests invalid as far as the exact determination of the flow stress is concerned. The large values of B indicate that the data should be friction corrected. However, the small number of tests does not provide sufficient information to extract friction data to perform an appropriate friction correction. Therefore, instead of stress–strain curves, mean pressure strain curves are given. Specimens deformed at 900 ◦ C at all the strain rates exhibited cracks along the circumference with the crack running in the direction of the compression axis. 4.2. Flow curves The recorded raw data were converted into mean pressure strain curves using the following relations. The strain ε is given by ε = ln
h , h0
Fig. 1. Flow curves for all experiments. The dotted line indicates the positions of peak mean pressure values for the determination of the activation energy.
presented in Fig. 2. As expected, the largest effect is observed for the highest strain rate. The flow curves can be temperature corrected if the activation energy for hot working QHW is known. 4.3. Determination of activation energy The activation energy for hot working was calculated using the method described in [13], which assumes that the influence of temperature T and strain rate ε˙ on flow stress σ is given by QHW A exp(βσ) = ε˙ exp , (5) RT where A and β are constants and R is the universal gas constant. Mean pressure values were determined by taking the intersection of the flow curve and the dotted line in Fig. 1. With these values an initial activation energy of 505 kJ mol−1 was estimated. However, as can be seen by comparing Figs. 1 and 2, these peak values are already affected by significant deformation heating. Therefore, the initial activation energy of 505 kJ mol−1 was used to correct the ex-
(3)
where h is the height during each stage of the test and h0 is the initial height. The mean pressure P is given by 4F , (4) πd 2 where F is the force at any given time and d is the corresponding specimen diameter at that time. The mean pressure strain curves obtained in this way are shown in Fig. 1. Although all curves exhibit a clear maximum mean pressure, it is debatable if this is a result of dynamic recrystallisation or deformation heating. From the recorded temperature data it is evident that deformation heating takes place, as can be seen from typical strain temperature curves for a nominal deformation temperature of 900 ◦ C, which are P=
Fig. 2. Strain temperature curves for a nominal deformation temperature 900 ◦ C and three strain rates.
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elling was also used to determine the hoop stress experienced for different test conditions, since it was found that the hoop stress for deformation at 900 ◦ C exceeded the fracture stress of the material, as evident by the occurrence of cracks. The
Fig. 3. Temperature-corrected flow curves. The dotted line indicates the positions of mean pressure values for the determination of the activation energy in last iteration. Experiment D is not shown due to the loss of thermocouple in the early stages of experiment. Without a temperature recording during the test, a temperature correction could not be performed.
perimental flow curves for deformation heating. From these temperature-corrected flow curves, a new activation energy was derived. The iteration process was repeated until no change in activation energy within the error boundaries was observed. The error was defined as the variation of activation energy in the flow stress region investigated here, which covers the range from 150 to 450 MPa. The final value for the activation energy by this method was found to be 487 ± 31 MPa. The temperature-corrected flow curves using this final activation energy are presented in Fig. 3. 4.4. Strain distribution obtained from FE model An example for the strain distribution obtained from the FE model is shown in Fig. 4, which represents deformation at a nominal temperature of 900 ◦ C and a strain rate of 10 s−1 . The positions labelled in this figure will be explained below in conjunction with microstructural investigations. FE mod-
Fig. 4. Strain distribution obtained from FE model for experiment C. Labels (a)–(d) mark positions, from which optical micrographs in Fig. 6(a)–(d) were taken.
Fig. 5. Microstructure of as-sprayed material (a) optical micrograph of polished sample demonstrating the presence of a carbide network, (b) optical micrograph of etched sample, (c) SEM image of etched sample.
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C. Rodenburg et al. / Materials Science and Engineering A 386 (2004) 420–427 Table 2 Summary of critical strain and strain rate at position where carbide network was broken down completely, obtained from FE modelling Experiment
Tnominal (◦ C)
εcrit
ε˙ (s−1 )
A C G
900 900 1100
0.88 0.89 0.82
0.18 16.8 0.17
minimum hoop stress for deformation at 900 ◦ C was obtained for a strain rate of 0.1 s−1 and was estimated to be 1.8 GPa. For deformation conditions resulting in smaller hoop stresses (i.e. at higher deformation temperatures), no cracks were observed. However, for deformation at 1000 ◦ C and strain rate of 10 s−1 with an estimated maximum hoop stress of 1.9 GPa, no cracks were observed. Therefore, the fracture stress of the investigated high-speed steel is assumed to be in the range 1.8–1.9 GPa over the given temperature range.
Fig. 6. Series of optical micrographs along compression axis of specimen deformed in experiment C. Positions corresponding to (a)–(d) are indicated in Fig. 4.
Fig. 7. SEM images of cross-sections perpendicular to compression axis of compression test samples (a) experiment B: nominal deformation temperature 900 ◦ C, (b) experiment H: nominal deformation temperature 1100 ◦ C. Note the difference in density of submicron-sized carbides and the presence of cracked carbides in (a).
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4.5. Influence of compression test parameters on microstructure The initial microstructure of the as-sprayed material is shown in Fig. 5. As far as the carbide distribution is concerned three different length scales were present: I. carbide network around 70 m consisting of irregular shaped carbides (highlighted by a black line in Fig. 5a), and close to elliptical or round carbides with average size of about 10 microns (marked by white lines in Fig. 5a); II. close to elliptical or round carbides of micron size (examples marked by white circles in Fig. 5c); III. rod shaped and rectangular carbides of sub micrometer size (Fig. 5c). Carbides of category I and II were visible by optical microscopy of polished samples (Fig. 5a). Carbides of group III were only visible in SEM images (Fig. 5c). However, a darkening in optical micrographs of the nital etched material (Fig. 5b) is indicative of their presence. This can be seen by comparing Fig. 5b and c. From these figures it is also evident that the carbide density varies from cell to cell. In order to investigate the influence of compression test parameters on microstructure, these different length scales have to be considered separately. The influence of strain on the carbide network structure is shown in Fig. 6, which shows a series of optical micrographs taken along the compression axis of the test specimen, as indicated in Fig. 4. No deformation of the cells is observed in Fig. 6a, taken from the area in contact with the tool during compression. Moving towards the centre of the specimen the carbide cell is first seen to be deformed in Fig. 6b, and then only local remnants are observed in Fig. 6c. There were no indications for the presence of a carbide network in the centre of the specimen (Fig. 6d). This behaviour was observed for all tests and indicates that the carbide network is broken down by deformation amounting to an equivalent plastic strain between 0.82 and 0.89, depending on deformation temperature. Table 2 summarises values obtained from FE modelling for the critical strain εcrit and strain rates ε˙ at which no carbide network remains. A strong influence of strain rate was not observed. A higher deformation temperature resulted in a slightly lower critical strain. The deformation temperature influences the carbide network in the direction perpendicular to the compression axis. The cracking of large irregular-shaped carbides was observed for deformation at 900 ◦ C (Fig. 7a) but not at 1100 ◦ C (Fig. 7b). Carbides falling into category II are not influenced by high stresses occurring at low deformation temperature, since their round shape and smaller size make them less prone to cracking [14]. Furthermore, these particles were identified by electron beam microanalysis as VC-type carbo-nitrides, which do not dissolve at the temperatures involved. The number of carbides of category III strongly depends on the deformation temperature, as is evident by comparing Fig. 7a and b. In these figures, the small carbides are rep-
Fig. 8. Optical micrographs for specimens deformed at strain rates given and nominal temperatures T = 900 ◦ C (a), T = 1000 ◦ C (b) and T = 1100 ◦ C (c).
resented by the small white dots in the low-magnification SEM images. The low magnification was chosen to provide a better impression of the density distribution over a large area, whereas the insets demonstrate the actual shape and size
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Table 3 Summary of activation energies for hot working a number of tools steels Steel
Q in kJ mol−1
Composition (wt.%) C
Cr
Mo
W
V
T1
0.95
4.35
0.11
18.41
0.99
AISI M3:2 M2 D2 H13 A2 W1
1.31 0.84 1.45 0.38 1.00 1.03
4.12 4.00 11.09 5.25 5.00 0.07
5.34 5.00 0.72 1.30 1.15 0.03
6.24 6.50 – – – –
3.05 1.90 0.71 1.05 0.3 –
of the type III carbides. The high density of carbides at the lower deformation temperature of 900 ◦ C in Fig. 7a is apparent. A low carbide density in conjunction with a strong local density variation was observed for deformation at 1100 ◦ C (Fig. 7b). The deformation temperature also influences the structure of the matrix, as can be seen from Fig. 8. For deformation below 1100 ◦ C (Fig. 8a and b), the matrix grain boundaries were not revealed by nital etching. However, following deformation at 1100 ◦ C, a well-developed matrix grain structure, resembling a recrystallised structure, was observed (Fig. 8c). However, recrystallisation did not seem to have taken place homogenously, and was only observed in the bands of lower carbide density (represented by absence of small black dots) in Fig. 8c.
5. Discussion The activation energy for hot working of 487± 31 kJ mol−1 is in the upper range of activation energies found for other tool steels, as can be seen from Table 3. However, it is conceivable that this is due to the high carbon content in combination with a very high vanadium content in the form of vanadium carbo-nitrides. A change in activation energy due to precipitation, as described in [5], was not observed in the temperature range tested here. The temperature-corrected flow curves for T < 1100 ◦ C do not exhibit maxima that would indicate the presence of dynamic recrystallisation, although dynamic recrystallisation is found to take place in other tool steels at this temperature and strain range for comparable strain rates [5,6]. The corresponding microstructures for T < 1100 ◦ C do not show signs of recrystallisation. In contrast, deformation at 1100 ◦ C does result in a recrystallised microstructure even though there is no pronounced peak in the flow curve. Therefore, static recrystallisation may have taken place during the time (about 1 s) between the end of deformation and the onset of quenching. Furthermore, within the high strain region recrystallisation seems to take place inhomogenously, confined to areas with a low density of submicron carbides. This is consis-
654 467 487 455 428 401 399 286
Ref. <1000 ◦ C >1000 ◦ C
[5] Current work [15] [16] [4] [15] [16]
tent with the presence of closely spaced carbides pinning the grain boundaries (Zener drag mechanism) and hindering recrystallisation [17]. The same mechanism is thought to be responsible for the absence of recrystallisation at deformation at 900 and 1000 ◦ C. As shown in Fig. 7, the carbide density of submicron-sized carbides is much higher at the lower deformation temperature. This higher volume fraction seems to be sufficient to prevent recrystallisation. One of the main aims of hot working of the investigated high-speed steel is the elimination of the carbide network. It was shown with the example of M2 that the dominant mechanism for breaking up the carbide network is mechanical fragmentation as opposed to a diffusion-controlled mechanism [7]. This is supported by this work, where Fig. 6 demonstrates progressive fragmentation with increasing strain. However, a small influence of deformation temperature on the critical strain (0.89 ≥ εcrit ≥ 0.82) was found (see Table 2), indicating that diffusion is involved to a small extent. In the current work, the critical strain for fragmentation of the carbide network was determined for a single deformation. However, in the normal forging of billets, the forging process would be undertaken in progressive smaller strain steps. Nevertheless, the total strain would allow not only break-up of the carbide structure but also static recrystallisation to refine the matrix structure and relax the residual strain in between deformations. The formation of cracks during hot working is undesirable. This work demonstrates that the hoop stress during deformation at 900 ◦ C exceeded the fracture stress and, therefore, should be avoided. Applying the values determined in these experiments to industrial practice would require a FE model of the particular industrial forming operation, highlighting the surface stress maxima.
6. Conclusions The hot workability of spray formed AISI M3:2 type high-speed steel was investigated in order to specify suitable forging conditions for this steel. From Arrhenius plots an activation energy for hot working of 487 ± 31 kJ mol−1 was determined. The combination of FE modelling and the
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observation of cracks for deformation at 900 ◦ C, but not at 1000 ◦ C, led to an estimated fracture stress in the range of 1.8–1.9 GPa. By observing microstructural changes as a function of specimen position and linking this to the strain distribution obtained by FE modelling, a critical strain, which is required to break up the carbide network, typically present in tool steels, was estimated to be 0.89 ≥ εcrit ≥ 0.82 covering the temperature range of 900–1100 ◦ C. This critical strain value has practical implications, for example the forging of spray formed high-speed steel billets into work rolls for hot rolling mills. It was found previously that the presence of the carbide network is responsible for increased oxidative wear at elevated temperatures [18]. Hence, the wear should be reduced if the forging conditions are tailored to exceed the critical strain obtained in this work.
Acknowledgements The authors would like to thank A. Lacey for performing the compression tests and E. Condliffe, School of Earth Sciences, University of Leeds for the electron microprobe analysis. This work was funded by EPSRC under grant number GR/R80209/01. The assistance of Danspray and Sheffield Forgemasters Rolls is also gratefully acknowledged.
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