1. Introduction

Recently high- and ultra high-strength low alloy TRIP-aided steel associated with transformation induced plastic-ity (TRIP)1) was developed for the automotive applicationsto attain weight reduction and to improve the impact safetyof vehicles. The TRIP-aided steel possesses excellentformability.2–9) Therefore it is expected to be used for theautomotive impact members and center pillars.

At present, TRIP-aided steels with polygonal ferrite2–7)

matrix (TDP steel), bainitic ferrite8) matrix (TBF steel) andannealed martensite9) matrix (TAM steel) have been devel-oped. Particularly, TBF steel is able to obtain a high tensilestrength more than 1 000 MPa. In this case, however, it is aconcern that delayed fracture occurs for the TBF steels inthe same way as conventional high strength steels.10–13) Theauthors14) reported that TBF steel has low hydrogen embrit-tlement susceptibility and high delayed fracture strength byhydrogen trapping on retained austenite. Since it is known

that retained austenite characteristics of TBF steel greatlyaffect delayed fracture properties, it is considered whetherthe addition of alloying element affects delayed fractureproperties of TBF steel.

In the present study, the effect of aluminum addition onhydrogen absorption behavior and delayed fracture proper-ties of the TBF steel were investigated. Also, they were re-lated with microstructural features and retained austenitecharacteristics of the steel.

2. Experimental Procedure

In this study, four kinds of vacuum melted and forgedslabs with different aluminum content (steels A through D)as listed in Table 1 were produced under condition of con-stant content of silicon and aluminum (Si�Al�1.5 mass%).For comparison, steel E with increased manganese contentwas also prepared. Martensite-start temperature (MS) of thesteels in the table was estimated by the following equa-

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ISIJ International, Vol. 48 (2008), No. 6, pp. 824–829

Effects of Aluminum on Delayed Fracture Properties of Ultra HighStrength Low Alloy TRIP-aided Steels

Tomohiko HOJO,1) Koh-ichi SUGIMOTO,2) Youichi MUKAI3) and Shushi IKEDA4)

1) Dept. Mech. Engng., Tsuyama National College of Technology, 624-1 Numa, Tsuyama 708-8509 Japan.2) Dept. Mech. Syst. Engng., Shinshu Univ., 4-17-1 Wakasato, Nagano 380-8553 Japan.3) Research & Development Labs., Kobe Steel Ltd., 2222-1 Ikeda Onoe-cho, Kakogawa 675-0023 Japan.4) Mater. Research Lab., Kobe Steel Ltd., 1-5-5 Takatsukadai, Nishi-ku, Kobe 651-2271 Japan.

(Received on January 16, 2008; accepted on March 31, 2008; originally published in Tetsu-to-Hagané,Vol. 93, 2007, No. 3, pp. 234–239 )

To improve the delayed fracture strength of ultra high-strength low alloy TRIP-aided steels with bainiticferrite matrix (TBF steels), the effects of aluminum content on hydrogen absorption behavior and delayedfracture properties of 0.2%C–0.2–1.5%Si–1.5%Mn TBF steel were investigated. When aluminum wasadded to the TBF steel, the diffusible hydrogen increased. It was expected that the hydrogen was chargednot only in retained austenite films but also on lath boundary. Delayed fracture strength of TBF steels con-taining aluminum were significantly increased, compared with conventional TBF steel. This was mainlycaused by (1) suppression of the stress-assisted martensite transformation resulting from the stabilized orcarbon-enriched retained austenite, (2) hydrogen trapping to refined interlath retained austenite films andlath boundary, and (3) relaxation of localized stress concentration by TRIP effect of the retained austenite.

KEY WORDS: TRIP-aided steel; ultra high strength steel; aluminum; retained austenite; hydrogen; delayedfracture.

Table 1. Chemical composition (mass%) and estimated martensite-start temperatures (MS, MS*, °C) of steels and retained austenitephases, respectively.

tion.15)

MS (°C)�550�361�(%C)�39�(%Mn)�0�(%Si)

�30�(%Al)�5�(%Mo) ............................(1)

where %C, %Mn, %Si, %Al and %Mo represent contentsof individual alloying elements, respectively.

Hot-rolling process and heat-treatment diagram are illus-trated in Fig. 1. First, slabs were heated to 1 200°C and hot-rolled to 3.2 mm thickness to a finishing rolling temperatureof 850°C and then cold-rolled to 1.2 mm thickness. Next,the sheet steels were annealed at 1 000°C for 1 200 s inaustenite region and then austempered at 300–475°C for200 s in salt baths to make the TBF steels with differenttensile strength.

Volume fraction of retained austenite was quantifiedfrom integrated intensity of (200)a , (211)a , (200)g , (220)gand (311)g peaks by X-ray diffractometry using Mo-Ka ra-diations.16) The carbon concentration (Cg mass%) was esti-mated from the following equation. In this case, lattice con-stant was measured from (200)g , (220)g and (311)g peaksof Cu-Ka radiation.17)

ag�3.5780�0.0330Cg�0.00095Mng�0.0056Alg

�0.0220Ng�0.0051Nbg�0.0031Mog ................(2)

where Mng, Alg, Ng, Nbg and Mog represent concentrationof individual elements (mass%) in retained austenite, re-spectively. In this study, we conveniently used added con-tent instead of these concentrations.

Tensile tests were carried out on a hard type of testingmachine at 25°C and at a crosshead speed of 1 mm/min(strain rate of 8.33�10�4/s), using JIS14B-type tensilespecimens of 15 mm gauge length, 6 mm width and 1.2 mm

thickness.Hydrogen was charged by the cathode charge method.

The charging bath composition and current density arelisted in Table 2. Total amount of solute hydrogen wasmeasured by TCD (Thermal Conduction Detector). Theamount of diffusible hydrogen was measured by TDS(Thermal Desorption Spectrometry). Specimen charged hy-drogen was kept in liquid nitrogen to prevent hydrogen evo-lution from the specimen.

Constant load tests by 4-point bending14) were conductedin a H2SO4�KSCN solution with hydrogen charging at25°C using rectangular specimens of 65 mm length, 10 mmwidth and 1.2 mm thickness. Maximum fracture strengthenduring for 5 h was defined as delayed fracture strength(DFL) in this study.

3. Results

3.1. Microstructure and Tensile Properties

Typical micrographs of TBF steels are shown in Fig. 2.As austempering treatment was carried out at temperaturesbelow MS, the matrix structure of all TBF steels consists ofbainitic ferrite and martensite with high dislocation density.Retained austenite films are located at these lath bound-aries. Addition of aluminum tends to refine these lath struc-tures and retained austenite films (Fig. 2(b)). On the otherhand, the microstructure of TBF steel with manganese isnot changed so much (Fig. 2(c)).

Tensile properties and retained austenite characteristicsof TBF steels are listed in Table 3. Tensile strength (TS) ofthese steels are between 1 108 and 1 420 MPa, and the totalelongation (TEl) are between 9.3 and 16.7%. Tensilestrengths of steels B, C and D decrease with increasing alu-minum content because of silicon removal of the same con-tent. On the other hand, manganese addition increases thetensile strength even when austempered at temperaturesabove 375°C. Yield ratio (YR) of the TBF steels are in-creased by aluminum and manganese addition. It is not ob-served that the addition of aluminum and manganese has aneffect on total elongation of TBF steels.

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Fig. 2. Typical image quality maps of steels (a) A, (b) D and (c) E austempered at 325°C, in which “gR” represents re-tained austenite.

Table 2. Hydrogen charging conditions.

Fig. 1. Hot and cold rolling and heat treatment process of TBFsteels.

Initial volume fractions of retained austenite (fg0) of steelA are between 3.4 and 4.8 vol%, and its carbon concentra-tion (Cg0) are between 0.85 and 1.30 mass%. Initial volumefractions of retained austenite of TBF steels with aluminumare the same grade as steel A. However, its carbon concen-tration increases with added amounts of aluminum. On theother hand, in steel E, initial volume fraction of retainedaustenite is increased, with a decrease in Cg0. These resultsare the same as previous studies.18,19)

3.2. Hydrogen Absorption Property

Figure 3 shows the relationship between total chargedhydrogen concentration (HT) and tensile strength (TS ) ofTBF steels after hydrogen-charging for 15 min. The totalcharged hydrogen concentration is hardly changed by alu-minum addition. When manganese is added, the totalcharged hydrogen concentration apparently increases.

Figure 4 shows hydrogen evolution curves of typicalTBF steels. The diffusible hydrogen increases with addedaluminum content, although total charged hydrogen con-centrations of steels A and D are almost the same, as shownin Fig. 4. In this case, the peak temperature of the evolutionrate of steel D is the same as the base steel (steel A). On theother hand, manganese addition raises the peak tempera-ture, with an increase in diffusible hydrogen concentration.

3.3. Delayed Fracture Strength

Figure 5 shows typical applied bending stress (sA)–timeto fracture (tf) curves of TBF steels. Figure 6 shows thevariations in delayed fracture strength (DFL) of TBF steelsas a function of tensile strength (TS). In TBF steels, delayedfracture occurs at more than 1 200 MPa. It is found that thedelayed fracture strength is increased by aluminum addi-tion, especially by aluminum addition of 0.5 mass%. On the

other hand, manganese addition considerably deterioratesthe delayed fracture strength.

Typical scanning electron micrographs of delayed frac-ture surface of TBF steels are shown in Fig. 7. Steels A andD exhibit mainly quasi-cleavage fractures. On the otherhand, intergranular fractures take place in steel E.

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Table 3. Tensile properties and retained austenite characteris-tics of steels used.

Fig. 3. Variations in total charged hydrogen concentration (HT)as a function of tensile strength (TS ) of steels A–E. Hy-drogen charging time: tC�15 min, and current density:500 A/m2.

Fig. 4. Comparison of hydrogen evolution curves of steels A, Dand E. Hydrogen charging time is tC�15 min and currentdensity is 500 A/m2.

Fig. 5. Typical applied bending stress (sA)–time to fracture (tf)curves of steels A, D and E with tensile strength of about1 300 MPa.

Fig. 6. Variations in delayed fracture strength (DFL) as a func-tion of tensile strength (TS ) in steels A–E.

4. Discussion

4.1. Effect of Alloying Elements on Hydrogen Absorp-tion Properties

In this study, a large amount of total charged hydrogenconcentration was found in steel E (Figs. 3 and 4). Gener-ally, hydrogen is trapped at carbide/matrix interface,20,21) ondislocation22) and on grain boundary21) in many ultra high-strength steels. The present TBF steels did not contain anycementites, although TBF steels included a large amount ofretained austenite films, particularly in steel E (5.5–11.9 vol%). According to a previous study,14) total chargedhydrogen concentration was linearly related with an initialvolume fraction of the retained austenite in 0.4C–1.5Si–1.5Mn TBF steel (980–1 960 MPa). As shown in Fig. 8, asimilar result was obtained in the present study. MoreoverChan et al.23) have reported that martensitic steels contain-ing retained austenite absorbed more hydrogen thanmartensitic steels without retained austenite, and this resultwas caused by a great deal of hydrogen trapping at retainedaustenite/martensite interface. Therefore, the high totalcharged hydrogen concentration of steel E may be causedby a high volume fraction of retained austenite. Also, it wasindicated that these hydrogen absorption properties were in-dependent of its carbon concentration. In this case, it wasexpected that most of hydrogen were trapped in retainedaustenite films and/or at retained austenite/matrix inter-faces.

In this study, in TBF steel containing manganese, thepeak temperature shifted to a higher temperature with in-creasing amount of diffusible hydrogen concentration. Onthe other hand, there was high dH/dt at the same peak tem-perature as steel A, although the amount of total chargedhydrogen concentration of steel D was hardly changed. Itwas reported by Tsubakino et al.11) that the hydrogen whichevolved at about 100°C corresponded to hydrogen trappingat grain boundaries, at carbide/matrix interfaces, on dislo-cations and on vacancies, and a peak of about 130°C washydrogen evolved from retained austenite. Furthermore itwas pointed out by Gu et al.24) that an increase of lathboundary area led to increase in the hydrogen trapping sitein 1 500 MPa grade bainite/martensite dual-phase steelscontaining retained austenite. The lath microstructure ofTBF steel was refined by aluminum addition, with refinedlath structure (Fig. 2). Therefore, it was considered that alarge amount of hydrogen evolution at high temperature ofsteel E was associated with the existence of a great deal ofretained austenite and an increase in dH/dt at peak tempera-

ture of about 110°C of the steel D was associated with in-crease in the lath boundary area.

4.2. Strain-induced Transformation and Stress-as-sisted Transformation of Retained Austenite

Generally, stability of retained austenite is deterioratedby hydrogen absorption. Figure 9 shows volume fraction ofretained austenite of TBF steels after delayed fracture tests.This figure suggested that volume fraction of retainedaustenite was decreased by hydrogen absorption, and as ap-plied stress was higher, volume fraction of retained austen-ite become lower. Also, deterioration of volume fraction ofretained austenite of steel D having high carbon concentra-tion of retained austenite was smaller than that of steel E.These results were able to be explained by the mechanismof martensite transformation and the change in free en-ergy25) (Fig. 10).

MS temperature of austenite (MS* in Table 1) of steel A(TA�325°C) calculated from Eq. (1) was more than roomtemperature. On the other hand, MS temperature of austen-

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Fig. 7. Scanning electron micrographs of fracture surface of steels (a) A, (b) D and (c) E austempered at 325°C.

Fig. 8. Relationship between total charged hydrogen concentra-tion (HT) and initial volume fraction of retained austenite(fg0) for steels A–E.

Fig. 9. Relationship between ratio of applied stress to yieldstress (sA/YS ) and volume fraction ratio ( fg /fg0) of steelsA, D and E.

ite of steel D (TA�325°C) was considerably low comparedwith delayed fracture testing temperature. However retainedaustenite became unstable because of hydrogen absorption.Therefore MS and MS

s temperatures were increased. Then itwas easy to change from retained austenite to martensitewhen a bending load was applied to the specimen (hydro-gen-assisted and/or stress-assisted transformation). How-ever, in steel D, the influence of increasing MS and MS

s tem-peratures by hydrogen absorption was removed by deterio-rating MS and MS

s temperatures by high carbon concentra-tion of retained austenite. So it was considered that thetransformation of retained austenite of steel D occurred notby hydrogen-assisted and/or stress-assisted but strain-in-duced transformation.

On the other hand, in steel A, though testing temperaturewas above MS temperature, transformation of retainedaustenite was assisted by hydrogen and/or stress. MoreoverT0 and MS temperatures of retained austenite seemed to in-crease because of effects of hydrogen and stress. This pro-moted hydrogen-assisted and/or stress-assisted transforma-tion of retained austenite with lower carbon concentration.

4.3. High Delayed Fracture Strength of TBF SteelContaining Aluminum

In this study, steels B, C and D achieved higher delayedfracture strength than steel A. Generally aluminum additionhardly changes the volume fraction of retained austenite,but promotes carbon-enrichment in retained austenite ifaustempering treatment is employed after annealing, be-cause aluminum anticipates the start of ferrite transforma-tion26) and suppresses cementite precipitation. Moreover,aluminum increases T0 temperature which inspires the con-centration of carbon to retained austenite during austemper-ing treatment.18) Resultantly, aluminum addition enhancedthe mechanical stability of retained austenite (and con-strained hydrogen-assisted and/or stress-assisted transfor-mation). In addition, yield stresses of the TBF steels wereincreased by aluminum addition (the yield ratios were alsoincreased), although tensile strength of the TBF steels withaluminum were deteriorated. These higher yield ratiosbring lower applied stress compared with yield stressenough. Therefore, it was considered that TBF steels con-taining aluminum were suppressed stress-assisted transfor-mation of retained austenite and increased delayed fracturestrength.

Generally, when retained austenite charged with much

hydrogen transforms to martensite, a lot of hydrogenevolves to other hydrogen trapping sites such as on grainboundary, on dislocation etc. because there are large differ-ences in hydrogen absorption properties between retainedaustenite and martensite. In this study, quasi cleavage frac-ture was observed on fracture surface of steels A and D,differing from the intergranular fracture of steel E, asshown in Fig. 7. In addition, TBF steel containing alu-minum indicated refined lath structure, and retained austen-ite existed refined and uniformly along the lath boundary.The lath boundary is able to become a valid hydrogen trap-ping site24) as well as retained austenite. So, it is consideredthat the hydrogen exists uniformly and hydrogen concentra-tion at the grain boundary was decreased in the steels con-taining aluminum because of hydrogen trapping on refinedlath boundary and in retained austenite. Moreover, in theTBF steels with aluminum, since hydrogen did not evolvefrom retained austenite by the suppression of hydrogen-as-sisted and/or stress-assisted transformation of retainedaustenite, hydrogen diffusion to other hydrogen trappingsites was suppressed. In addition, the localized stress con-centration was relaxed by strain induced plasticity of stabi-lized retained austenite. This may contribute to suppresscrack and/or void initiations at the grain boundary.

5. Conclusions

Hydrogen absorption properties and delayed fracturestrength of TBF steels containing aluminum were investi-gated to improve the delayed fracture resistance of TRIP-aided steels. The results are summarized as follows.

(1) The total charged hydrogen concentration washardly influenced by aluminum content. However, hydrogenevolution concentration of TBF steel with aluminum in-creased, compared with base steel without aluminum, al-though the peak evolution temperature was nearly the sameas that of steel A. This was because the hydrogen wascharged on lath boundary which was refined by the alu-minum addition.

(2) Delayed fracture strength of TBF steels containingaluminum of 0.2–1.0 mass% increased vastly, compared tothe TBF steel without aluminum. This result was explainedas follows.(i) Hydrogen-assisted and/or stress-assisted martensite

transformation of retained austenite was suppressedby aluminum addition, because the retained austenite

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Fig. 10. Influence of temperature on (a) martensite transformation mechanisms and (b) free energy change.

was stabilized or carbon-enriched.(ii) The diffusible hydrogen of the TBF steel containing

aluminum was trapped on lath boundary increased byrefining due to aluminum addition and in refined andstabilized retained austenite.

(iii) The localized stress concentration was relaxed bystrain induced plasticity of stabilized retained austen-ite.

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