ARCHAEOMATERIALS

Damascus Steel Revisited

J.D. VERHOEVEN ,1,5 A.H. PENDRAY,2 W.E. DAUKSCH,3

and S.R. WAGSTAFF4

1.—Iowa State University, Ames, IA, USA. 2.—Williston, FL, USA. 3.—Nucor Corp, Charlotte,NC, USA. 4.—Novelis Inc, Sierre, Switzerland. 5.—e-mail: jver@iastate.edu

A review is given of the work we presented in the 1990s that successfullydeveloped a technique for reproducing the surface patterns and internalmicrostructure of genuine Damascus steel blades. That work showed that a keyfactor in making these blades was the addition of quite small levels of carbide-forming elements, notably V. Experiments are presented for blades made fromslow- and fast-cooled ingots, and the results support our previous hypothesisthat the internal banded microstructure results from microsegregation of Vbetween dendrites during ingot solidification. A hypothetical model was pre-sented for the mechanism causing the unique internal microstructure that givesrise to the surface pattern forming during the forging of the ingots from whichthe blades are made. This article attempts to explain the model more clearly andpresents some literature data that offer support to the model. It also discusses analternate model recently proposed by Foll.

INTRODUCTION

The Damascus steel swords found mainly inmuseums became famous because of their attractivesurface patterns and their reputation originating inthe crusades as being sharper and tougher thanequivalent western swords. However, there was amystery about how these swords were made becauseafter around 1850 no one could make one. C.S.Smith, the chief metallurgist in the Manhattanproject, brought the mystery of these blades to theattention of the metallurgy community in a bookpublished in 1960.1 He described what was knownabout the metallurgy of these swords and how theywere made from small wootz steel ingots producedin India. It was known that the Damascus swordswere forged from small hockey puck-sized high-carbon steel ingots made in India that are calledwootz steel in the west. Smith presented micro-graphs of the internal microstructure of a bladeshowing alternating sheets of pearlitic steel andsheets of clustered small spheroidized carbide par-ticles, Fe3C. This laminated internal microstructure

produces the surface patterns on the swords aftersuitable etching. It is largely this unique surfacepattern that has made the Damascus sword sofamous. C.S. Smith and other participants at a 1985conference on Damascus steel blades held at NewYork University2 agreed that the art of makingthese blades had been lost shortly after around theearly nineteenth century.

In the 1990s we carried out an extensive set ofexperiments3–5 in which bladesmith Alfred Pendraywas able to produce knives that matched both theinternal microstructure and external patterns ofgenuine Damascus steel swords. The experimentsproduced small hockey puck-sized ingots of steelshaving carbon compositions similar to those ofgenuine Damascus steel swords of around 1.5–1.7wt.% C. The aligned cluster sheets of small well-spheroidized carbide particles occurred on forgingthe ingots to blade shapes only if they contained asmall amount, around 0.01%, of carbide-formingelements, such as V, Mo or Cr, with V being mosteffective. The result was consistent with the factthat six genuine museum-quality Damascus swordswere also found to contain V at low levels.3 Ingotsforged to blade shapes of around 4–5-mm thick-nesses produced microstructures and surface pat-terns closely matching those of museum-qualityblades.

A.H. Pendray: Deceased.

JOM, Vol. 70, No. 7, 2018

https://doi.org/10.1007/s11837-018-2915-z� 2018 The Minerals, Metals & Materials Society

(Published online May 10, 2018) 1331

CONTROLLING THE SPACINGOF THE CLUSTER SHEETS

The blades consist of alternating sheets of clus-tered cementite particles and sheets of pearlite. Thepearlite sheets can contain cementite particles. Asthe number of cementite particles in the pearlitesheets is reduced, the overall sheet sharpness isimproved. When viewing the surface pattern of aDamascus blade from a comfortable distance, theboldness of the pattern perceived by the eyeimproves with sharper sheet formation for a givensheet spacing. Our experience examining bothmuseum blades and blades made by Pendray havefound that for a given sheet sharpness, the surfaceboldness is maximized with sheet spacings ofaround 40–100 microns. Our previous work3–5

hypothesized that the microstructure is formed bya type of banding due to the microsegregation of Vbetween the dendrites during the solidification ofthe ingots. One would therefore expect the spacingto be set by either the primary or the secondarydendrite spacing in the ingot from which a blade isforged combined with the amount the ingot isreduced on forging. If controlled by the primaryspacing, the resulting cluster sheet spacing wouldbe larger than if it were controlled by a combinationof secondary and primary spacings. Our initialstudies used a phosphorus addition of around 0.1wt.% to match the composition measured on severalwell-patterned museum-quality blades. At this levelof P, the ingots are hot short and very difficult toforge. We think this is why a 1722 report byReaumur claimed that Parisian blacksmiths wereunable to forge the small wootz cakes from India.1

The P forms a ternary eutectic called steadite,which becomes molten at the forging temperatures,thereby causing the embrittlement.

To overcome this problem we found that it wasnecessary to heat the ingots packed in an iron millscale for around 10–12 h at 1100–1200�C. Thistreatment produces a rim of nearly pure iron onthe ingot and is termed a rim heat treatment. Theiron rim adequately contains the molten steaditewithin the ingot so that it can be forged withoutcracking. Apparently, the Parisian bladesmiths hadnot learned the trick of the rim heat treatment.(Pendray was able to obtain an ancient partiallyforged bar of steel verified to be a Damascus barfrom the Anwar Armory in India. Additional forgingof this bar produced a typical surface Damascuspattern, and metallographic sections of the barverified that it had been given a rim heat treat-ment6). Eliminating the P addition removed the hotshortness problem. The blades still produced excel-lent cluster sheet formation but, unfortunately,without the rim heat treatment the cluster sheetspacing was found to be significantly smaller. Onsuch blades forged to thicknesses of 3–5 mm, thespacings were running in the 20–30 micron range,and to easily see the surface pattern, it was

necessary to use a magnifying glass. During therim heat treatment, diffusion of V between thesmaller distances would occur, which would beneeded to homogenize the V compositions oversecondary spacings. At the same time there wouldbe some level of microsegregation of V between theprimary spacings.

Experiments were recently carried out to produceblades made from ingots prepared as previouslydescribed3,4 and cooled at slow and fast coolingrates. The fast rate was obtained by turning off thepropane gas to the furnace that heated the enclosedcrucibles. The slow rate was obtained by a slowreduction of propane gas. Figure 1 presents thetime–temperature history of the ingots made at thetwo cooling rates. The temperature was measuredat the bottom outside edge of the crucible with aplatinum thermocouple (type B). Experiments withthermocouples placed both below the bottom of thecrucible and in a ceramic protection tube inside thecrucible at the bottom of the melt showed thatoutside temperatures ran 33 ± 4�C higher and18 ± 4�C lower than the inside melt temperatureon heating versus cooling, respectively. Chemicalanalysis of ingot 8115 is presented in Table I. Theiron carbon phase diagram7 gives a liquidus tem-perature of 1423�C for 1.57% C. Hence, the melt washeated to a maximum temperature of 1468�C andheld above the liquidus of 1423�C (� 1390 on heatup and � 1441 on cool down on the plot) for � 65min. It is interesting to point out that when cruciblesteel was being produced in the late nineteenthcentury, to sufficiently kill the steel and avoidporosity in the ingots, they simply held the meltuntil it stopped bubbling.8,9 They had to hold themelt temperature above the freezing temperaturefor 30–40 min before gas evolution stopped. Theauthors do not specify what actual temperature wasused. It must be that times of at least 30–40 min arerequired for the [O] + [Fe] M CO equilibrium to beachieved adequately to avoid cavity porosity ontheir cooled ingots. Apparently the 30-min hold ofingot 8115 adequately killed the steel so that cavityporosity did not occur at the cooling rate of thisingot. The cooling rate given by the slope of thecurve during subsequent dendrite growth wasaround 4.3�C/min.

Such was not the case at slower cooling rates.Several attempts were made to cool ingots at slowerrates, but large cavity porosity was found in theingots, which made it nearly impossible to forgethem to blade shapes. It seemed most likely thatthis porosity was due to evolution of CO gas oncooling at the slow rates. Cavity porosity in theslow-cooled ingot 7315 was avoided by adding aborosilicate glass ampoule filled with 1.4 g of Al and1.4 g of a chloride-based fluxing salt. The glassampoule plus the fluxing agent prevented the Alfrom oxidizing before it could dissolve into the meltand thereby prevented CO evolution just as it doesin aluminum-killed steels. Table I shows that the

Verhoeven, Pendray, Dauksch, and Wagstaff1332

carbon composition of ingot 7315 is the same as thatof ingot 8115; hence, it has the same liquidustemperature of 1423�C. It was heated to a maximumtemperature of 1468�C, held above the liquidus for108 min and subsequently cooled through the den-drite growth range at an average rate of around 1�C/min. There was no cavity porosity in this slow-cooled ingot.

Both ingots were given a rim heat treatment of16 h at 1040�C. The ingots were 90 mm diameter by43 mm height. Both ingots were forged to a bladeshape with dimensions of 43 mm wide by a bitunder 8 mm thick. A total of 36 forging cycles were

employed with the first 7 carried out over a tem-perature range of � 1030�C to 700�C and theremainder over the range of � 925�C to 700�Cfollowed by three thermal cycles between 950�Cand 700�C. The blades were forged to a thickness of7.75 ± 0.05 mm, and both displayed excellent clus-ter sheet alignment. As expected, the slower cooledingots produced blades with a larger average clustersheet spacing, 93 lm versus 71 lm for ingot 8115.These results present additional evidence that theformation of the aligned cluster sheet spacing is aresult of the microsegregation of the carbide-form-ing element V between the dendrites during solid-ification of the ingot. Blade 8115 was forged furtherto a thickness similar to a knife blade or a sword,2.8 mm, and the cluster sheet spacing dropped to 45microns.

Figure 2 presents a micrograph of blade 7315 at7.8 mm thickness, and the enlargement illustratesthe nature of the clustered arrays of cementiteparticles in the cluster sheets as well as the pearlitesheets between the cluster sheets. The amount ofcementite in the pearlite sheets has been reducedadequately to produce a bold surface pattern.

DISCUSSION OF PROPOSED MODELFOR PATTERN FORMATION

As discussed in the supplementary material, thedendritic solidification of steel results in the forma-tion of sheets of dendrites separated by interden-dritic sheets. Previous work presented ahypothetical model4,5 for the mechanism of theformation of the laminated internal microstructurethat forms on forging. The model will be more fullyexplained here and some literature data that offersupport to the model presented. An alternate modelthat has come to our attention will also bediscussed.

In a previous study,10 metallographic examina-tions of the evolution of the microstructure offorged ingots at various stages of the forging

Fig. 1. Temperature versus time for ingots 8115 and 7315.

Table I. Chemical analysis of fabricated blades

Element 7315 8115

C 1.57 1.57Mn 0 0V 0.061 0.047Mo 0.01 0.01Cr 0.02 0.02Zr 0.002 0Nb 0.009 0.002Ti 0.003 0.001S 0.006 0.008Sn 0.009 0.003P 0.02 0.025Al 0.014 0.008B 0.0008 0.0003Cu 0.01 0.01Ca 0.0002 0.0007Ni 0.02 0.02Pb 0.0082 0.0007N < 0.016 < 0.016

Fig. 2. Longitudinal section of a blade from ingot 7315.

Damascus Steel Revisited 1333

sequence from ingot shape to thin blade shapehave been evaluated, and micrographs are shownin Fig. 3. It was found that the large blockycementite carbides formed in the original ingotwere quickly replaced by small spheroidizedcementite particles after initial forging at hightemperatures. We think the formation of the smallspheroidized carbide particles was enhanced by theforging flow because as the temperature dropsbelow the Acm temperature, sheet-like carbidesformed on austenite grain boundaries would breakup because of the forging flow and quickly spher-oidize at the high temperatures. The arrangementof the carbide particles in the forged bars wasfound to be random after the 90-mm-diameter by43-mm-high ingot had been forged into a bladeshape with a 36-mm-wide by 18-mm thick crosssection. This took 27 forging cycles. The carbideparticles appear to have nucleated on prior austen-ite grain boundaries as demonstrated by the triplejunction shown at the arrowhead in Fig. 3a as wellas on prior austenite grain boundaries. The carbidedistribution at this stage of forging appears fairlyrandom. On additional forging, it was found thatthe cementite carbide particles slowly began toarrange themselves into the cluster sheet geometrywith the density of clustering and the alignment ofthe sheets improving as the number of forgingcycles increased. After 70 forging cycles when theblade thickness was reduced to 5 mm, Fig. 3c, thecluster sheet geometry became well formed. So, thecritical question is what is causing the carbides tocluster together into the aligned cluster sheets inthe later stages of forging.

DISCUSSION OF THE PATTERN FORMA-TION MODEL AND SUPPORTING LITERA-

TURE DATA

The proposed model requires a selective coarsen-ing during the multiple forging cycles of the cemen-tite particles in the interdendritic sheets where the

V has been concentrated. In pure iron/carbon steels,cementite has the chemical formula Fe3C. Whenalloying elements dissolve into the cementite, it iscustomary to designate the chemical formula asM3C where M = Fe + X, with X being primarily V inthis case. During the heat-up stage of a forgingcycle, the cementite particles will begin to dissolveinto the neighboring austenite. If the highest tem-perature does not exceed the Ac-cm temperature,some of the cementite particles will remain at thehighest temperature and will then grow in size asthe temperature falls. If the alloying element in theM3C carbide reduces the mobility of the carbide/austenite interface, then the M3C carbides in theinterdendritic sheets will become larger than theFe3C carbides in the dendritic sheets as the tem-perature rises. There are three possibilities:

Case 1 At the highest forging temperature, bothsome Fe3C and some M3C particles survive. Butthe M3C particles are larger and have a highernumber density.

Case 2 At the highest forging temperature, all ofthe Fe3C particles have been dissolved but someof the M3C particles have survived.

Case 3 At the highest forging temperature, all ofthe cementite particles have dissolved.

On cool down, the cementite particles will grow insize. In case 1, it might be possible for the remainingFe3C particles, which grow faster than the M3Cparticles, to increase their size and match that ofthe larger but slower growing M3C particles. Then,if enough of the Fe3C particles nucleate in thedendritic sheets to match the number density in theinterdendritic sheets, no enhanced coarseningoccurs in the interdendritic sheets. But if few orno additional Fe3C carbide particles nucleate in thedendritic sheets on cool down, one would find larger

Fig. 3. Longitudinal sections of a blade after forging to sizes of (a) 36 9 18 mm, 27 cycles; (b) 36 9 14 mm, 34 cycles; and (c) 45 9 5 mm, 70cycles.

Verhoeven, Pendray, Dauksch, and Wagstaff1334

carbides in the interdendritic sheets after each cooldown step of the forging, which eventually wouldcause the selective coarsening of carbides in theinterdendritic sheets as the number of forging cyclesincreases.

Nucleation of new cementite particles is a rela-tively difficult process requiring some supercooling.If the carbon being forced out of the austenite as thetemperature drops can deposit onto existing car-bides it will do so rather than nucleating additionalcarbides. The M3C carbides in the interdendriticsheets will be close by the dendritic sheets, and it ispostulated that the carbon will simply diffuse to theM3C carbides rather than nucleating new Fe3Ccarbides in the dendritic sheets. The net effect willbe the selective coarsening of cementite particles inthe interdendritic sheets that is observed in exper-iments such as shown in Fig. 3.

Case 2 would produce the selective coarseningeven more quickly during the forging cycles becausethe number density of Fe3C carbides in the dendriticsheet region would begin at zero at the start of cooldown. Case 3 will be considered later.

If the hypothesized mechanism for coarsening iscorrect, it requires that:

Requirement 1 Cementite carbides in interden-dritic sheets contain a significant amount of Vdissolved into them, while cementite carbides in thedendritic sheets ae V free.

Requirement 2 The dissolved V must reduce themobility of the austenite/cementite interface duringgrowth or shrinkage processes.

Requirement 1 As explained in the supplementarymaterial, an electron microprobe study5 found thatcarbides formed in the interdendritic sheets con-tained significant amounts of dissolved V, whilethose in dendritic sheets were V free.

Requirement 2 As shown in the supplementarymaterial data on the effect of the carbide formingelements, Mo, Cr and Mn on hardness loss duringtempering [S3] offer strong experimental supportthat dissolved carbide particles significantly reducethe mobility of cementite/iron interfaces.

AN ALTERNATE POSSIBLE MECHANISMFOR CLUSTER SHEET FORMATION DURING

FORGING

H. Foll11 has hypothesized an alternate mecha-nism for cluster sheet formation during forging thatcould occur for case 3 discussed above. The carbide-forming elements form native carbides such as VCfor V. These native carbides form at higher temper-atures than Fe3C and would be stable at the highestforging temperatures; hence, they could act asnucleating agents for cementite formation in thecooling step of forging. This hypothesis is discussedin detail in the supplementary section, and severalreasons are presented for why it is unlikely. One of

the main reasons is that although native carbideswere found in the interdendritic regions of V and Moblades in the study of Ref. 5, no native carbides werefound in the Cr and Mn steels, which showed clustersheet formation.

CONCLUSION

The pattern on genuine Damascus steel swordsand blades is produced by the formation of alignedsheets of clustered cementite particles separated bysheets of pearlite. Experiments are presented onblades made from ingots solidified at slow and fastcooling rates, and the results are consistent withour previous hypothesis that this microstructure isproduced by a type of banding caused by microseg-regation of V between the dendrites formed duringsolidification. Some previous work10 is reviewed,which show that the formation of the carbide clustersheets requires the presence of very low levels ofcarbide-forming elements and only occurs aftermany forging cycles and that the alignment andsharpness of the cluster sheets improve with thenumber of forging cycles. The hypothetical model forthe mechanism of this effect involving selectivecoarsening of the cementite carbides in interden-dritic sheets presented by the authors4, 5 has beenreviewed in some detail. An alternate model sug-gested by H. Foll involving nucleation of cementiteon primary carbide has been shown to be possiblebut unlikely. In addition, literature data on theeffect of carbide-forming elements on reducinghardness loss of steels on tempering after quenchingis shown to offer support for the model.

ACKNOWLEDGEMENTS

The chemical analyses presented in Table I werecarried out by Mark Schmidt at Nucor Steel Mill,Darlington, SC. The metallography was done withequipment provided by Iowa State University, and theglass ampule was prepared with equipment at MIT.

ELECTRONIC SUPPLEMENTARYMATERIAL

The online version of this article (https://doi.org/10.1007/s11837-018-2915-z) contains supplemen-tary material, which is available to authorizedusers.

REFERENCES

1. C.S. Smith, A History of Metallography, Chapters 3 and 4(MIT Press, Cambridge MA, 1960 and 1988).

2. D.W. Ecker and G.N. Pant, The Damascus Blade: Legendsand Realities. in Proceedings of Symposium On the Blade-smiths Art (New York University, New York, NY 1985).

3. J.D. Verhoeven, A.H. Pendray, and W.E. Dauksch, J. Met.50, 58 (1998).

4. J.D. Verhoeven, A.H. Pendray, and E.D. Gibson, Mat.Char. 37, 9 (1996).

5. J.D. Verhoeven, A.H. Pendray, and W.E. Dauksch, Trans.ISS Iron Steelmak. 25, 65 (1998).

Damascus Steel Revisited 1335

6. J.D. Verhoeven, A.H. Pendray, andW.E.Dauksch, J.Met. 56, 17(2004).

7. Metals Handbook, 8th edn. vol. 8, (American Society forMetals, 1973), p. 275.

8. O.M. Becker, High Speed Steels (New York: McGraw-Hill,1910), p. 58.

9. Making Shaping and Treating of Steels, 7th edn. (U.S.Steel Co., 1957), p. 264.

10. J.D. Verhoeven and A.H. Pendray, Mat. Char. 30, 175 (1993).11. H. Foll, 11.5 Wootz Swords (2017), https://www.tf.

uni-kiel.de/matwis/amat/iss/kap_b/backbone/rb_5_1.html#_dum_2.

Verhoeven, Pendray, Dauksch, and Wagstaff1336