METALLURGY

High dislocation density–inducedlarge ductility in deformedand partitioned steelsB. B. He,1 B. Hu,2 H. W. Yen,3 G. J. Cheng,3 Z. K. Wang,4 H. W. Luo,2* M. X. Huang1*

A wide variety of industrial applications require materials with high strength and ductility.Unfortunately, the strategies for increasing material strength, such as processing tocreate line defects (dislocations), tend to decrease ductility.We developed a strategy tocircumvent this in inexpensive, medium manganese steel. Cold rolling followed by low-temperature tempering developed steel with metastable austenite grains embedded in a highlydislocated martensite matrix.This deformed and partitioned (D and P) process produceddislocation hardening but retained high ductility, both through the glide of intensive mobiledislocations and by allowing us to control martensitic transformation.The D and P strategyshould apply to any other alloy with deformation-induced martensitic transformation andprovides a pathway for the development of high-strength, high-ductility materials.

Strength and ductility are key mechanicalproperties of metallic materials for de-veloping energy-efficient and lightweightstructural components in a wide variety ofindustries, including automotive and aero-

space. Unfortunately, improving strength oftenresults in the degradation of ductility, which isknown as the strength-ductility trade-off (1). Costis also an issue, as alloying elements that helpimprove both properties, such as cobalt and ti-tanium, tend to be expensive. Previous effortstoward resolving this trade-off focused on engi-neering defects such as grain boundaries (2, 3)and coherent twin boundaries (4, 5). However,the strength may reach a limit when grain sizeor twin-boundary spacing is reduced to nano-meters (6–8). Dislocations or line defects are adifferent pathway for engineering alloy proper-ties. In contrast to the above planar defects, thestrength of metallic materials monotonically in-creases with dislocation density according tothe well-known Taylor hardening law (9). Theproblem is that ductility tends to decrease asthe number of dislocations increases. However,improved choices in alloy composition and clev-er processing strategies may allow dislocationengineering to circumvent the strength-ductilitytrade-off. In particular, achieving high ductilityin the presence of high dislocation density in aninexpensive material produced by facile process-ing routes is desirable for broad industrial appli-cations at an economic cost.Severe plastic deformation (SPD) and phase

transformation both generate high dislocation

density in metals. Cold drawing is one SPD thatintroduces a very high dislocation density inpearlitic steel wires (10). Cold drawing and otherSPDmethods also may lead to grain refinement(<100 nm) in metals because of the formationof high-angle grain boundaries by rearrange-ment of intensive dislocations (10, 11). By contrast,the first-order solid-state martensitic transforma-tion produces a highly dislocated martensite mi-crostructure without a lot of grain refinement(~500 nm) in steels (12). Either quenching ordeformation can trigger martensitic transforma-tion, depending on the alloying content of steels(13, 14).We show that rolling and low-temperaturetempering produced a high dislocation densityin steel, also enabling a large ductility. In ad-dition to the high dislocation density, intersti-tial C atoms partition into and help stabilize theaustenite phase during tempering (13). We definethis thermomechanical treatment as the “deformedand partitioned (D and P)” process and the cor-responding steel as D and P steel. The D and Pprocess could be realized with conventionalprocessing techniques that are compatible withexisting industrial production lines.We used a medium Mn steel (10% Mn, 0.47%

C, 2% Al, and 0.7% V, by weight) for the D and Pprocess. The Mn and C atoms are effective aus-tenite stabilizers (fig. S1). The addition of 2% Alcontent is to suppress cementite precipitationduring the tempering process. The addition of0.7% V is to form intensive nanometer-sized Vcarbides, which provided enhanced resistanceto delayed fracture induced by hydrogen em-brittlement (15, 16). The D and P steel that weproduced bymultiple deformation and annealingsteps (17) possesses a heterogeneous lamelladual-phase microstructure in which metastableaustenite is embedded in a martensite matrix(Fig. 1A). We obtained themartensite matrix bycold rolling (fig. S2) followed by tempering. Thetempered martensite matrix has heterogeneousgrainmorphologies and substructures. Large len-

ticular martensite grains, which constitute mostof the matrix, are mainly decorated with disloca-tion cells (Fig. 1B). Nevertheless, some large len-ticularmartensite grains also possess dislocationsas well as twins (Fig. 1C) and some small lathmartensite grains with dislocations (Fig. 1D) inthe martensite matrix. We estimated the averagedislocation density of martensite matrix (~1.27 ×1016m–2) with synchrotron x-ray diffraction (XRD)(fig. S3). Such a high dislocation density is dueto the combination of plastic deformation anddisplacive shear transformation. The martensitewe created differs from the thermally formedmar-tensite obtained by quenching in that it inheritsdislocations generated from warm deformation.The early transformed martensite is also sub-jected to cold rolling, which we confirmed bythe formation of dislocation cells in martensite(Fig. 1B) (18). Consequently, the dislocation den-sity in our martensite matrix is much larger thanthat in thermally formed martensite (~2.36 ×1014 m–2) (fig. S3). The austenite phase also hasa heterogeneousmicrostructure, including coarselamella grains (Fig. 1A), ultrafine lamella grains(Fig. 1D), and submicron granular grains (Fig. 1E).The large austenite grains are decorated with ahigh number of dislocations as well as stackingfaults (Fig. 1F), whereas the small austenite grainsmainly possess dislocations (Fig. 1E). Both mar-tensite and austenite phases contain nanometer-sized vanadium carbides (Fig. 1, G and H) andsolid solutes (C,Mn, andVatoms). Although somelayered austenite grains have higher Mn content(fig. S4) from the banding of Mn in hot-rolledstrip (fig. S1), we also observed austenite grainswith low Mn content and martensite with highMn content (fig. S4). This suggests that the Mnbanding does not determine retention of thelayered austenite grains in the D and P steel.Instead, it could be due to the mechanical sta-bilization induced by intensive defects in theselayered austenite grains (Fig. 1F) (19).The as-developed D and P steels possessed

ultrahigh yield strengths of 2.21 and 2.05 GPa(Fig. 2). The high yield strength results from thehigh dislocation density (rt) in the martensitematrix. For instance, the yield stress (s) thatresults from dislocation forest strengthening isestimated to be 1.6 GPa according to theTaylor hardening law (9)

s ¼ Mambffiffiffiffi

rtp

where Taylor factorM is 2.9 (20), empirical con-stant a is 0.23 for cell-forming dislocation dis-tribution (21), shear modulus m is 85 GPa (22),and Burgers vector b is 0.25 nm (20). Therefore,the yield stress induced by the high dislocationdensity is on the same order of magnitude as themacroscopic yield strength (Fig. 2). The collectivecontributions from the other strengthening de-fects, such as solid solutes, nanoprecipitates,stacking faults, twin boundaries, and ultrafinegrains (Fig. 1), elevated the yield strength beyond2 GPa. Consequently, the D and P steel has ahigher yield strength than steel with a bimodalcoarse-grainedmicrostructure (Fig. 2 and fig. S5).

RESEARCH

He et al., Science 357, 1029–1032 (2017) 8 September 2017 1 of 4

1Department of Mechanical Engineering, The University of HongKong, Pokfulam Road, Hong Kong, China. 2School ofMetallurgical and Ecological Engineering, University of Scienceand Technology Beijing, Xue Yuan Lu 30, Beijing 100083, China.3Department of Materials Science and Engineering, NationalTaiwan University, Roosevelt Road, Taipei 10617, Taiwan.4Department of Mechanical and Biomedical Engineering, CityUniversity of Hong Kong, Tat Chee Avenue, Hong Kong, China.*Corresponding author. Email: mxhuang@hku.hk (M.X.H.);luohaiwen@ustb.edu.cn (H.W.L.)

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Intriguingly, the D and P steel also had a largertensile uniform elongation (>15%) than its bi-modal coarse-grained counterpart. Our D and Psteel with high dislocation density has an inversestrength-ductility trade-off similar to what hasbeen found in other materials (23).The high dislocation density in our D and P

steel results in an ultrahigh yield strength. The

large uniform elongation is also due to the dis-location density. The D and P steel demonstratesdiscontinuous yielding followed by large Lüdersstrains (e) (Fig. 2). The enhanced yield drop afterthe prolonged annealing is a typical feature ofstrain aging during which both mobile and im-mobile dislocations are segregated by solid solutes(24). Martensite has an inherently high ratio

(~18%) of mobile screw dislocations from dis-placive shear transformations (25, 26). Themobileand immobile dislocations in martensite couldrearrange to form dislocation cells (Fig. 1B) dur-ing cold rolling, while the characteristics ofdislocations remain the same (18, 27). The dis-location cells consist of a cell interior with a lowdislocation density and a cell wall with a highdislocation density. The catastrophic release andglide of the mobile screw dislocations from thecell wall lead to the collapse and break of orig-inal cell boundaries and the formation of elon-gateddislocation cell structure after Lüders strain,which we confirmed by comparing Fig. 1B andFig. 3A. The elongated cell has an increasedlength of ~400nmcompared to that of the originalcell structure, which indicates that the averageglide distance of the mobile screw dislocationsis ~400 nm. A general relationship between plasticstrain and mobile dislocation density (rm) canbe written as (28)

e ¼ rmbL=M

where L is the average glide distance of themobile dislocations (~400 nm). We assumed

He et al., Science 357, 1029–1032 (2017) 8 September 2017 2 of 4

Fig. 2. Engineering stress-strain curves ofour steels at room temperature. Curves a andb represent the D and P steels, processed bywarm rolling, intercritical annealing, and coldrolling followed by tempering at 400oC for 6 and15 min, respectively. Curve c represents thedeformed steel, processed by warm rolling,intercritical annealing, and cold rolling withouttempering at 400oC. Curve d represents thesample with bimodal coarse-grained micro-structure (fig. S5) obtained by annealing at900oC for 30 min. The uniform elongation isalmost the same as the total elongation for alltensile curves. An increased yield strength of180 MPa in curve a as compared to that in curve c is due to the segregation of C atoms at dislocationsduring the tempering at 400oC (24).The serrations at elongations between 10 and 14% in curves a andb manifest the Portevin-Le Châtelier effect (35).

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Fig. 1. Microstructureof our D and P steelbefore the tensile test.(A) Electron backscatterdiffraction phase imageand orientation imageshowing the lamellamicrostructure of austen-ite (g) embedded in tem-pered martensite (a′)matrix. The austenite areafraction is 15%, and themartensite area fractionis 85%. RD indicatesthe rolling direction, NDindicates the normaldirection, and TD indi-cates the transversedirection. (B) Dislocationstructures in martensite.The inset is a selected-area diffraction pattern(SADP). (C) A typicallenticular martensite dis-tributed with twins anddislocations. (D) The lathmartensite and lamellaaustenite as observed intransmission electronmicroscopy (TEM) bright-field and dark-fieldimages. The right imageis a SADP showing aKurdjumov-Sachs (K-S)relation between lamellaaustenite and lath mar-tensite. (E) Dislocations in submicron granular austenite. (F) Dislocations and stacking faults in a large austenite grain captured by TEM bright-fieldand high-resolution images. (G) Distribution of nanosized vanadium carbide (V4C3) in tempered martensite matrix. (H) High-resolution TEM image ofnanosized vanadium carbide in austenite.

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the ratio of mobile dislocation density to totaldislocation density to be 18% for deformation-induced martensite (26), given a mobile disloca-tion density of 2.3 × 1015 m–2 in the martensitematrix. Consequently, the plastic strain inducedby the glide of these mobile dislocations in themartensite matrix is estimated to be 6.8%, whichconstitutes a large portion of measured Lüdersstrain (~7 to 9%) in the D and P steel (Fig. 2). Thehigh mobile dislocation density in the martensitematrix accommodates a large plasticity uponyielding. The catastrophic release and glide ofthe mobile screw dislocations result in a low, butpositive, strain-hardening rate during Lüders strainregime (fig. S6). Such a low but positive strain-hardening rate has been reported in irradiatedmetals (29) and w-containing Ti alloys (30), inwhich the operation of dislocation channelingis the governing deformation mechanism.In addition to the high mobile dislocation

density, the D and P steel had a continuoustransformation-induced plasticity (TRIP) effectat a large strain during the tensile test (Fig. 3B).The TRIP effect resulted from the formation ofmartensite in the coarse layered austenite grains.This TRIP effect applied compressive residualstress to effectively blunt localizedplasticity duringtensile straining (fig. S7) (31) and also provideddynamic strain partitioning between phases andimproved strain hardening (fig. S6) (32). Theformation of the largemartensite grains suggeststhat coarse layered austenite grains are less stablethan ultrafine austenite grains during uniaxialtensile deformation (Fig. 3C). Nevertheless, mostof these coarse layered austenite grains onlytransform to martensite at strains beyond theLüders strain (Fig. 3B), explaining their highmechanical stability in the D and P steel. Weascribed the enhancement in the mechanicalstability to the high dislocation density in largeaustenite grains (Fig. 1F), where dislocations canact as barriers for glissile martensite interfaceand therefore stabilize austenite grains (19).More-over, the hard martensite matrix (fig. S8) canshield austenite from deformation (33), allowingthe austenite to transform at a large strain re-gime. The austenite grains in the D and P steelwere further stabilized by C partitioning frommartensite (fig. S9) (13) relative to the austenitegrains in deformed steel (fig. S10). In addition tothe effect of C partitioning and relatively higherMn content (fig. S4), the high dislocation densityin the D and P steel also controlled release of theTRIP effect, improving ductility.The D and P steel also had a twinning-induced

plasticity (TWIP) effect during the tensile test.We mostly observed the TWIP effect, inducedby the formation of deformation twins, in ul-trafine austenite grains (Fig. 3D). The initiationof nanotwins from phase boundary suggests ahigh stress level experienced by lamella aus-tenite grain (Fig. 3D). Therefore, the TWIP ef-fect also operated in the large strain regime(fig. S6). The minor austenite volume fraction(Fig. 1A) of the small grains means that we ex-pect the TWIP effect to be much less importantthan the TRIP effect, even as deformation twins

He et al., Science 357, 1029–1032 (2017) 8 September 2017 3 of 4

Fig. 3. Microstructure of our D and P steel after the tensile test. (A) The elongateddislocation cell structure in the tempered martensite matrix after tensile straining to 8%. Theinset is the corresponding SADP. Arrows mark the break of cell boundaries. (B) The XRD profilesat different strains. The XRD measurement on the specimen of 5.9% strain is on the gaugepart swept by Lüders band. Compared to the (220)g peak, the substantially enhanced (211)a′peak intensity at strains beyond 5.9% suggests that martensitic transformation is active in thelarge-strain regime. a.u. represents arbitrary units. (C) The formation of lenticular martensitein the coarse austenite grains after tensile straining to fracture. (D) The generation ofdeformation twins in the submicron austenite grains after tensile straining to fracture. The rightimage is the SADP showing a K-S relation between twinned lamella austenite and lath martensite.

D&P Steel

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High Specific Stength SteelHigh Entropy AlloyNanostructured Lamella Ti

Nanostructured Mo AlloyNanostructured Co Alloy

Ti6Al4V

TungstenInconel718

d in Figure 2

a & b in Figure 2

Fig. 4. Tensile properties of our steel compared with those of other existing high-strengthmetallic materials. These include nanobainite steel (36), maraging steel (37, 38), dual-phasesteel (39), martensitic steel (40), TRIP steel (41, 42), quenching and partitioning (Q and P)steel (43), nanotwinned steel (44, 45), maraging-TRIP steel (46), high–specific strength steel(47), high-entropy alloys (48, 49), nanostructured lamella Ti (3), Ti6Al4V (47), tungsten (50),Inconel 718 (51), nanostructured Mo alloy (52), and nanostructured Co alloy (53). Our D andP steel is clearly separated from the general trend and therefore defines a new space in thestrength-ductility map. Note that the uniform elongation is selected here not only because it isa desirable property but also because it is less affected by the sample dimensions used indifferent studies (2).

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can accumulate deformations and improve duc-tility (4).We compared the bulk properties of D and P

steel to other high-strength metallic materials(Fig. 4). TheD and P steel exhibits a yield strengththat is 50% higher than that of nanobainite steelwhile maintaining a comparable uniform elon-gation (Fig. 4). Moreover, it exhibits a uniformelongation that is one order of magnitude largerthan that of commercial maraging steels whilemaintaining an equivalent yield strength. Con-sequently, our D and P steel achieves excellenttensile properties and defines a new space in thestrength-ductility map (Fig. 4). Despite its dis-continuous yielding and Lüders strain (Fig. 2),the ultrahigh yield strength of our D and P steelmakes it a desirable alloy for applications whereyield strength is the main design criterion.The D and P overcomes the challenge of

creating martensite, which is known to be anissue for compositionally similar steels that relyon a quenching and partitioning process (13)(fig. S5). The high dislocation density in the Dand P steel not only increases the yield strengthby dislocation forest hardening but also enablesa large ductility by the glide of existing mobiledislocations and by the controlled release of theTRIP effect. The high dislocation density is theorigin of the inverse strength-ductility trade-off.We expect that this strategywill be useful in othersystems with similar deformation-induced mar-tensitic transformation mechanisms such as tita-nium alloys (34). The D and P steel exhibits a lowraw-materials cost as compared to the maragingsteel while maintaining a comparable ultimatetensile strength (fig. S11). Therefore, by engi-neering dislocations, we simultaneously alleviatethe economic concernswhile achieving ultrahighstrength.

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ACKNOWLEDGMENTS

The authors thank K. Lu for his insightful and constructivecomments on this paper. M.X.H. is grateful for financial supportfrom the Research Grants Council of Hong Kong (grant nos.712713, 17203014, and 17255016) and the National NaturalScience Foundation of China (grant no. U1560204). H.W.L. isgrateful for financial support from the National NaturalScience Foundation of China (grant no. U1460203). H.W.Y.acknowledges the Ministry of Science and Technology of theRepublic of China for providing financial support undercontract MOST-104-2218-E-002-022-MY3 and thanks theInstrumentation Center at National Taiwan University fortechnical support during use of the JSM 7800F PRIME high-resolution scanning electron microscope. The authors alsoacknowledge the Shanghai Synchrotron Radiation Facility forproviding the synchrotron XRD facility at beamline no. 14 B. Thepresent work has pending patents with application numbers201610455155.3 and PCT/CN2016/096509 and anotherawarded patent, number 201410669029.9. M.X.H. supervisedthe study. M.X.H., B.B.H., and H.W.L. designed the study. B.B.H.and B.H. prepared the thermomechanical treatment and themechanical tests. H.W.Y., G.J.C., and B.B.H. conducted themicrostructure characterization. B.B.H., M.X.H., H.W.L., and Z.K.W.analyzed the data. B.B.H. and M.X.H. wrote the paper. Theauthors declare no competing financial interests. Data areavailable in the manuscript and supplementary materials.

SUPPLEMENTARY MATERIALS

www.sciencemag.org/content/357/6355/1029/suppl/DC1Materials and MethodsSupplementary TextFigs. S1 to S14References (54–58)

20 February 2017; accepted 10 August 2017Published online 24 August 201710.1126/science.aan0177

He et al., Science 357, 1029–1032 (2017) 8 September 2017 4 of 4

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induced large ductility in deformed and partitioned steels−High dislocation densityB. B. He, B. Hu, H. W. Yen, G. J. Cheng, Z. K. Wang, H. W. Luo and M. X. Huang

originally published online August 24, 2017DOI: 10.1126/science.aan0177 (6355), 1029-1032.357Science 

, this issue p. 1029Sciencelow-temperature annealing and contained a dislocation network that improved both strength and ductility.''forest'' of line defects in manganese steel. This deformed and partitioned steel was produced by cold-rolling and

used a processing mechanism to create aet al.the microstructure that increases strength tends to reduce ductility. He Many industrial applications require materials to have high strength while remaining pliable, or ductile. However,

A ductile steel shows its strength

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