Influence of the holding time of the deep cryogenic treatment on the strain hardening behavior of HY-TUF steel
© Zare et al. 2015
Received: 2 August 2015
Accepted: 9 November 2015
Published: 14 November 2015
The effect of the holding time of the deep cryogenic treatment on the strain hardening of HY-TUF, a low-alloy medium-carbon steel, and a comparison with strain hardening of conventional heat-treated steel were investigated.
For this purpose, a deep cryogenic treatment was done for different holding times of 12, 24, 48, and 72 h. The metallographic results indicate completion of martensitic transformation; η-carbide precipitation and decrement of carbon concentration in martensitic matrix happened while carrying out a deep cryogenic process.
Observations of transmission electron microscopy indicate higher density of dislocations within martensite near the cementite compared to that near the η-carbides. The tensile flow stress data for this steel was examined in terms of Hollomon equation. The results show that strain hardening of conventional heat-treated steel takes place in two stages.
This evidence is related to the co-deformation of austenite and martensite. The strain hardening takes place in one stage for the deep cryogenic-treated steels, and their strain hardening exponents increase (from ~0.29 to 0.47) with an increasing holding time up to 48 h. A further increase in the holding time of the deep cryogenic treatment is found not to vary strain hardening exponent. The increase of strain hardening exponent and then observation of plateau in this parameter show an optimum value for the holding time of the deep cryogenic treatment (48 h).
KeywordsStrain hardening Deep cryogenic treatment Martensite Austenite η-carbide
where n is the strain hardening exponent, ε the true strain, K the strength coefficient, and σ the true stress. The constant parameters of strain hardening in Eq. 1 consisting of K and n can be measured by drawing the true stress–strain data in a diagram on double natural logarithmic scale and fitting a straight line at it. The intercept at the unit value of the true strain obtains the K magnitude, while the slope of the line gives the n value. It is well known that the strain hardening parameters, K and n, are of significant industrial attention. The magnitude of n maintains quantification of the capability to decelerate localization of the plastic deformation of the materials. Materials with a high magnitude of strain hardening exponent are selected for operations, which involve plastic deformation because the higher the magnitude of n, the greater is the rate at which the material strain hardens (Meyers and Chawla 1991; Akbarpour and Ekrami 2008). The value of K provides the strength level of material or the magnitude of required forces in plastic deformation (Akbarpour and Ekrami 2008; Hertzberg 1996). Moreover, the unique material strength properties, ultimate tensile strength (UTS) and yield strength (YS), are individually important to consider and control as they influence the behavior of structures during accident in airplanes (Marin 1962). Taken together as the UTS/YS ratio (known as the “strain hardening value” in European practice) and strength difference (UTS-YS), they indicate the ductility capacity of the structural components where they were used (Fridman 1975). It is mentioned that the UTS/YS and UTS-YS can be considered as criteria for strain hardening.
HY-TUF steel is a specific class of low-alloy medium-carbon steel that derives its strength from martensitic transformation and precipitation of carbides in a quenched and tempered condition. In HY-TUF steel, compared to other low-alloy medium-carbon steels (such as AISI 4340 steel), C content is maintained at a low level (<0.26 wt%); the contents of Si and Mn are relatively high (~3 wt%), as well as Ni (~1.80 wt%). This design of the composition is considered for HY-TUF steel to get superior toughness and resistance to crack propagation as well as ultra-high strength. In addition, vacuum arc remelting (VAR) is used in the steel-making process of HY-TUF steel to have a clean material with low content of inclusions. HY-TUF steel has aerospace and aircraft applications where lightweight structure with ultra-high strength and toughness are essential and cost is not a major concern such as main landing gear, ultra-sensitive screws, and flexible drive shaft of helicopter (Klopp 1992).
Deep cryogenic treatment is a beneficial way in improving the mechanical properties of steels. Deep cryogenic treatment increases hardness, strength, and toughness and improves abrasive and fatigue wear resistance. Deep cryogenic treatment has been investigated to enhance mechanical properties of steels because it improves the microstructure. This process decreases or eliminates retained austenite (Leskovšek et al. 2006). In addition, a deep cryogenic process with post-tempering treatment causes precipitation of fine η-carbides from supersaturated martensite. These fine carbides have spherical morphology and a high homogenized distribution (Senthilkumar et al. 2011). The martensite lath is originated during deep cryogenic treatment, distributed more uniformly, and has smaller colonies, because the martensitic transformation is suppressed with the formation of fine η-carbides, which are coherent with the matrix (Li et al. 2010a, b).
There are generally two atomic mechanisms that describe the formation of η-carbide during a deep cryogenic process. The first mechanism is based on moving of carbon atoms towards the dislocations during long holding at cryogenic temperatures. The carbon clusters adjacent to the dislocations which are considered as nuclei for the η-carbide (Li et al. 2010a, b, 2011). Tyshchenko et al. (2010) considered the second hypothesis to interpret the carbon redistribution in the martensitic matrix. The suggested mechanism which was studied by using Mossbauer spectroscopy and internal friction measurements attributed the formation of η-carbide to capture locked carbon atoms by mobile dislocations. At present, little is understood about the stability of cryo-treated η-carbide which is a transient type. Recently, Hadi (2014) presented data of differential scanning calorimetry about aging of η-carbide in the cryo-treated HY-TUF steel. It is mentioned that η-type carbide is stable during aging at a temperature range of 200–400 °C and does not transform into more stable carbides such as cementite. This evidence is attributed to the coherency of interphase boundaries of η-carbide and martensite that was observed by a high-resolution transmission electron microscope.
The literature studies showed that even though the micromechanism behind the strain hardening and its relationship with deep cryogenic-treated (DCT) microstructure have not been clarified, different microstructural observations have been studied in the literatures (Li et al. 2010a, b; Akhbarizadeh et al. 2009; Amini et al. 2012; Chi et al. 2010; Jaswin and Lal 2011; Li et al. 2010a, b; Koneshlou et al. 2011). So the present research is undertaken to investigate the influence of the holding time of the deep cryogenic treatment on the strain hardening behavior of HY-TUF steel. The micromechanisms of strain hardening were realized by explaining the strain hardening behavior in terms of Hollomon equation. The strain hardening of conventional heat-treated (CHT) and DCT steels are also compared. In all cycles, the austenitizing and tempering parameters were considered constant to evaluate the effect of the holding time of the deep cryogenic treatment on the strain hardening phenomenon.
Chemical composition of the steel (wt%)
Tensile samples were cut from the forged bars of HY-TUF steel by a wire cut machine. The strain hardening behavior was determined using tensile testing at room temperature. Tensile tests were done according to ASTM E8M-04 standard with an Instron tensile testing machine at a cross-head speed of 2 mm min−1. The average strength of the three measurements was reported for each sample. In order to confirm the strength variation of martensite phase with cryogenic treatment, microhardness measurements were carried out using a Leitz RZD microhardness tester (Vickers indenter with a load of 150-g force).
Result and discussion
The typical TEM microstructures of the CHT and DCT48 steels are shown in Fig. 7. Before preparing TEM foils, the samples were not plastically deformed and just cryo-treated; thus, martensite as well as austenite phase have low density of dislocations. The dislocation density in the martensite phase adjacent to the η-carbide particles is low (Fig. 7b). Fig. 7a shows cementite particles which are located in the martensitic matrix. It is observed from this micrograph that the martensitic areas adjacent to the cementite involve high density of dislocations, unlike the martensite which is located close to the particles of η-carbides. Furthermore, the size of cementite (~320 nm) in CHT steel becomes much larger in comparison to η-carbides (~210 nm) in DCT48 microstructure.
Strain hardening behavior
According to Fig. 8, the DCT steels show typical linear behavior of strain hardening. Linear variation of ln σ with ln ε indicates that the investigated DCT steels, with low volume fractions of retained austenite (<1 %), obey the one-stage strain hardening mechanism. At this stage, martensite deforms plastically and η-carbide particles could not deform because of their complex crystalline structure (orthorhombic). It should be mentioned that volume fractions of austenite are low enough (<1 %) (Figs. 4 and 5); this evidence has no significant effect on the plastic deformation of the DCT steels (Sakaki et al. 1990).
The variation of strain hardening exponent
As can also be seen from Fig. 8 for the CHT sample, the first stage has a high strain hardening exponent (n 1) (~0.54) and the second stage (n 2) has a low one (~0.21). The n 1 is a criterion for strain hardening of austenite, while n 2 is related to the co-strain hardening of austenite and martensite. According to more ductile nature of austenite in comparison with that of martensite, it is expected that the first stage has a greater strain hardening exponent compared to that of the second stage.
Change in strain hardening exponent with holding time of the deep cryogenic treatment is also shown in Fig. 9. The peaks/plateaus trend (~0.47) is observed in the variation of this parameter as a function of holding time at the deep cryogenic temperature. The variation of strain hardening exponent can be interpreted as well as explained for changes of UTS/YS and UTS-YS in the “Result and discussion” section to the “The variations of UTS/YS and UTS-YS” section.
The variation of strength coefficient
It can also be seen in Fig. 10 that the strength coefficient of strain hardening stage 2 for the CHT steel is higher than that of stage 1. This shows that at the second stage, a phase with a higher level of yield strength (martensite) is strain hardening in comparison to austenite phase which is deforming at the first stage of strain hardening. It should be mentioned that K 1 of the CHT steel is only affected by the flow behavior of austenite while K 2 is affected by two parameters which are as follows: (i) austenite deformation can be ongoing at strain hardening stage 2, which leads to decrement in K 2 of CHT steel; and (ii) besides, as well as listed in Fig. 6, the existence of high-strength martensite in CHT steel and its deformation at the second stage of strain hardening, it compensates the decrement of the K 2 as a result of austenite plastic flow. Thus, this evidence is leading to a greater strength coefficient of stage 2 when compared to that of stage 1.
As can be observed in Fig. 10, with the increasing holding time of the cryogenic process up to 48 h, strength coefficient decreases and shows a plateau (~5100 MPa). During deep cryogenic process, η-carbides precipitate with the mechanism of stress-assisted diffusion, so carbon content of martensite decreases and as a consequence its strength decreases (Totten and Howes 1997). So, with the increasing holding time up to 48 h, more η-carbide nucleus forms and more carbon atoms are absorbed from martensitic matrix, and this evidence leads to decrement in strength coefficient. In holding times more than the optimum (72 h), the driving force for subzero diffusion (transformation and thermal stresses) decreases, and this leads to the invariable state of strength coefficient (Amini et al. 2012).
The variations of UTS/YS and UTS-YS
Increasing trend of the UTS/YS and UTS-YS with the holding time of the deep cryogenic treatment up to 48 h is also shown in Fig. 11. The peaks/plateaus trend of these parameters has been identified as one of the manifestation of η-carbide formation. With increasing holding time up to the peak value (48 h), more carbon atoms diffuse to form η-carbides. So with the formation of η-carbides in martensitic matrix, carbon content of martensite decreases and its stretch formability increases, which means enhancement in UTS/YS and UTS-YS as well as the strain hardening exponent (n) (Totten and Howes 1997). Also, with an increasing holding time more than 48 h, the UTS/YS and UTS-YS remained unaltered. It revealed that after 48 h, the internal stresses decrease to diffuse carbon atoms and unaltered trend in the UTS/YS and UTS-YS parameters (Amini et al. 2012).
Completion of martensitic transformation and η-carbides precipitation occur during cryogenic treatment. An increment in the degree of holding time of the deep cryogenic treatment up to 48 h significantly increases the volume fraction of the η-carbides and their uniformity while decreases the particle size of η-carbides. In longer holding durations more than 48 h, the volume fraction, uniformity, and particle size of carbides do not change anymore in deep cryo-treated steels. Precipitation of homogenized η-carbides within DCT martensite causes decrement in carbon concentration of the martensite as compared to that in the CHT steel. The results also show that exhaustion of carbon atoms from martensitic matrix and the precipitation of η-carbide are increased with the increasing holding time of the deep cryogenic treatment up to 48 h.
Using the diffraction patterns of TEM, the existence of η-carbides is confirmed in the deep cryo-treated steels. TEM micrograph of the CHT steel indicates high density of dislocations within the interphase boundaries of cementite/martensite, while the deep cryo-treated sample for 48 h contains η-carbides within low-carbon martensitic matrix. Additionally, η-carbide which forms during deep cryogenic treatment has homogenized spherical morphology as compared to the carbide of the CHT sample.
Strain hardening of CHT steel comprises two stages with different exponents. The higher strain hardening exponent (~0.54) at the first stage is attributed to the deformation of the austenite, and the lower strain hardening exponent (~0.21) at the second stage is associated with the plastic deformation of the both austenite and martensite.
The DCT steels show one-stage strain hardening behavior. At this stage of strain hardening, martensite phase deforms plastically.
By increasing the holding time of the deep cryogenic treatment, the strain hardening exponent shows peaks/plateaus (~0.47), while strength coefficient has a decreasing plateau trend. The variations of Hollomon equation parameters with holding time at the cryogenic temperature are related to the microstructural changes including precipitation of η-carbides and decrement in carbon content of martensite phase during cryogenic treatment.
The variation of Hollomon’ strain hardening exponent is confirmed by the changes of UTS/YS ratio and UTS-YS as the criteria for strain hardening.
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